High-strength steel sheet, hot-dipped steel sheet, and alloy hot-dipped steel sheet that have excellent fatigue, elongation, and collision characteristics, and manufacturing method for said steel sheets

ABSTRACT

This high-strength steel sheet includes: in terms of percent by mass, 0.03 to 0.10% of C; 0.01 to 1.5% of Si; 1.0 to 2.5% of Mn; 0.1% or less of P; 0.02% or less of S; 0.01 to 1.2% of Al; 0.06 to 0.15% of Ti; and 0.01% or less of N; and contains as the balance, iron and inevitable impurities, wherein a tensile strength is in a range of 590 MPa or more, and a ratio between the tensile strength and a yield strength is in a range of 0.80 or more, a microstructure includes bainite at an area ratio of 40% or more and the balance being either one or both of ferrite and martensite, a density of Ti(C,N) precipitates having sizes of 10 nm or smaller is in a range of 10 10  precipitates/mm 3  or more, and a ratio (Hvs/Hvc) of a hardness (Hvs) at a depth of 10 μm from a surface to a hardness (Hvc) at a center of a sheet thickness is in a range of 0.85 or more.

TECHNICAL FIELD

The present invention relates to a high-strength steel sheet, ahot-dipped steel sheet, and an alloyed hot-dipped steel sheet which aresteel sheets for automobiles and are mainly subjected to press working.In particular, the present invention relates to a high-strength steelsheet, a hot-dipped steel sheet, an alloyed hot-dipped steel sheet, andproduction methods thereof, and these steel sheets have excellentfatigue properties and excellent collision properties with a sheetthickness of about 6.0 mm or less and a tensile strength of 590 MPa ormore.

The present application claims priority on Japanese Patent ApplicationNo. 2009-127340 filed on May 27, 2009, the content of which isincorporated herein by reference.

BACKGROUND ART

In recent years, for the purpose of reducing weight and enhancing safetyof an automobile, an increase in the strength of automobile componentsand materials used therein has been made, and with regard to steelsheets which are representative materials for the automobile components,a rate of use of a high-strength steel sheet has been increased. Inorder to achieve the reduction in weight while enhancing safety, it isnecessary to increase a collision energy absorbing ability whileincreasing the strength. For example, it is effective to increase ayield stress of a steel material; and thereby, a collision energy can beabsorbed efficiently with a low deformation amount. In particular, as amaterial used in the vicinity of a cabin of an automobile, materialshaving high yield stresses are widely used because there is a need toblock a colliding object invading the cabin from the point of view ofoccupant protection. Particularly, the demand for a high-strength steelsheet having a tensile strength in a range of 590 MPa or more, and ahigh-strength steel sheet having a tensile strength in a range of 780MPa or more has been increasing.

In general, as methods of increasing a yield stress, there are (1) amethod of work-hardening a steel sheet by performing cold rolling, (2) amethod of forming a microstructure including a low-temperaturetransformation phase (bainite or martensite) having a high dislocationdensity as a main phase, (3) a method of performing precipitationstrengthening by adding microalloying elements, and (4) a method ofadding solid-solution strengthening elements such as Si and the like.Among them, with regard to the methods (1) and (2), the dislocationdensity in the microstructure is increased; and thereby, workabilityduring press forming is deteriorated drastically. This results infurther deterioration of press formability of a high-strength steelsheet which originally has insufficient in workability. On the otherhand, in the method (4) of performing solid-solution strengthening, theabsolute value of a strengthening amount is limited; and therefore, itis difficult to increase the yield strength to a sufficient extent.Accordingly, in order to efficiently increase the yield stress whileobtaining high workability, it is preferable that microalloying elementssuch as Nb, Ti, Mo, and V are added to perform precipitationstrengthening of alloy carbonitrides for achieving a high yield stress.

From the above viewpoint, a high-strength hot-rolled steel sheet inwhich precipitation strengthening of microalloying elements is utilizedhas been put to practical use. However, the high-strength hot-rolledsteel sheet in which the precipitation strengthening is utilized mainlyhas two problems. One is fatigue properties and the other is rustprevention.

With regard to the fatigue properties as the first problem, in thehigh-strength hot-rolled steel sheet in which precipitationstrengthening is utilized, there is a phenomenon in which a fatiguestrength is reduced due to softening of the surface layer of the steelsheet. In the surface of the steel sheet which directly comes intocontact with a rolling roll during hot rolling, the temperature of onlythe surface of the steel sheet is reduced due to a heat releasing effectof the roll which comes into contact with the steel sheet. When thetemperature of the outermost layer of the steel sheet falls below an Ar₃point, coarsening of the microstructure and precipitates occur; andthereby, the outermost layer of the steel sheet is softened. This is themain factor of the deterioration of the fatigue strength. In general, afatigue strength of a steel material is increased as the outermost layerof the steel sheet is hardened. Therefore, in a high-tensile hot-rolledsteel sheet in which precipitation strengthening is utilized, it isdifficult to obtain a high fatigue strength at present. On the otherhand, the purpose of increasing the strength of a steel sheet is toreduce the weight of an automobile body; however, the sheet thicknesscannot be reduced in the case where the fatigue strength ratio isreduced while the strength of the steel sheet is increased. From thispoint of view, it is preferable that the fatigue strength ratio be in arange of 0.45 or more, and even in the hot-rolled high-tensile steelsheet, it is preferable that the tensile strength and the fatiguestrength be maintained at high values with a good balance. Here, thefatigue strength ratio is a value obtained by dividing the fatiguestrength of a steel sheet by the tensile strength. In general, there isa tendency that a fatigue strength increases as a tensile strengthincreases. However, in a material with higher strength, the fatiguestrength ratio is reduced. Therefore, even though a steel sheet having ahigh tensile strength is used, since the fatigue strength is notincreased, there may be a case where a reduction in the weight of theautomobile body which is the purpose of increasing strength cannot berealized.

The other problem is rust prevention. Typically, as a steel sheet usedin a chassis frame for an automobile, a cold-rolled steel sheet producedby cold rolling and annealing thereafter and an alloyed hot-dipgalvanized steel sheet are not used, but a hot-rolled steel sheet havinga relatively thick thickness in a range of 2.0 mm or more is mainlyused. In the vicinity of a chassis where a paint on the surface of thesteel sheet is easily peeled off due to physical contact with curbs,flying stones, or the like, a material having a thicker thickness thanthat required from a design stress is selected to be used inconsideration of a corrosion thickness reduction amount (amount ofreduced sheet thickness due to corrosion) during a service life; andthereby, the quality is guaranteed. Therefore, with regard to thechassis frame and the like, the reduction in weight by substituting thematerial to a high-strength steel sheet is delayed at present, comparedto body components. Since the sheet thickness is thick as one of thecharacteristics of chassis components, arc welding is mainly conductedfor welding the components. Since the arc welding has a higher heatinput amount than that of spot welding, HAZ softening is more likely tooccur. In order to obtain properties of being resistant to HAZsoftening, precipitation strengthening by an addition of microalloyingelements is mainly utilized. Therefore, it is difficult to apply ahot-dip galvanized steel sheet or an alloyed hot-dip galvanized steelsheet having high rust prevention properties because annealing isconducted after cold rolling for the purpose of structure strengtheningin the manufacture of these galvanized steel sheets. The reason that theprecipitation strengthening by an addition of microalloying elementscannot be utilized for the steel sheet produced by performing annealingafter cold rolling is described as follows. Even in the case where ahot-rolled steel sheet into which microalloying elements are added issubjected to a cold rolling at a high cold rolling rate (for example,30% or higher) and then annealing is conducted at a temperature in arange of an A₃ point or less, the microalloying elements suppressrecovery and recrystallization of ferrite. Therefore, a microstructureis work-hardened in a state of being cold-rolled; and as a result,workability is deteriorated drastically. On the other hand, in the casewhere heating is performed at a temperature in a range of the A₃ pointor higher, precipitates coarsen; and as a result, there is a problem inthat a sufficient increase in the yield strength is not obtained.Therefore, the precipitation strengthening by the addition ofmicroalloying elements cannot be utilized.

As a hot-dip galvanized steel sheet which includes a hot-rolled steelsheet, Patent Document 1 discloses a method of producing a hot-dipgalvanized steel sheet having a tensile strength in a range of 38 to 50kgf/mm². With regard to the steel sheet having such a strength level, adesired strength level is obtained without utilizing precipitationstrengthening due to an addition of microalloying elements. However,methods of producing a high-strength steel sheet, a hot-dipped steelsheet, and an alloyed hot-dipped steel sheet, which have excellentcollision properties and fatigue strength with a strength in a strengthlevel of 590 MPa or more are not disclosed yet.

PRIOR ART DOCUMENT Patent Document

-   Patent Document 1: Japanese Examined Patent Application, Publication    No. H06-35647

DISCLOSURE OF THE INVENTION Problems to be Solved by the Invention

In order to solve the above-described problems, the present inventionaims to provide a high-strength steel sheet, a hot-dipped steel sheet,an alloyed hot-dipped steel sheet, and production methods thereof, andthese steel sheets have a tensile strength in a range of 590 MPa ormore, and are excellent in fatigue properties, elongation, and collisionproperties.

Means for Solving the Problems

The high-strength steel sheet of the present invention having excellentfatigue properties, elongation and collision properties, includes: interms of percent by mass, 0.03 to 0.10% of C; 0.01 to 1.5% of Si; 1.0 to2.5% of Mn; 0.1% or less of P; 0.02% or less of S; 0.01 to 1.2% of Al;0.06 to 0.15% of Ti; and 0.01% or less of N; and contains as thebalance, iron and inevitable impurities. A tensile strength is in arange of 590 MPa or more, and a ratio between the tensile strength and ayield strength is in a range of 0.80 or more. A microstructure includesbainite at an area ratio of 40% or more and the balance being either oneor both of ferrite and martensite. A density of Ti(C,N) precipitateshaving sizes of 10 nm or smaller is in a range of 10¹⁰ precipitates/mm³or more. A ratio (Hvs/Hvc) of a hardness (Hvs) at a depth of 10 μm froma surface to a hardness (Hvc) at a center of a sheet thickness is in arange of 0.85 or more.

In the high-strength steel sheet of the present invention havingexcellent fatigue properties, elongation and collision properties, afatigue strength ratio may be in a range of 0.45 or more.

An average dislocation density may be in a range of 1×10¹⁴ m⁻² or less.The high-strength steel sheet may further include one or more selectedfrom the group consisting of: in terms of percent by mass, 0.005 to 0.1%of Nb; 0.005 to 0.2% of Mo; 0.005 to 0.2% of V; 0.0005 to 0.005% of Ca;0.0005 to 0.005% of Mg; 0.0005 to 0.005% of B; 0.005 to 1% of Cr; 0.005to 1% of Cu; and 0.005 to 1% Ni.

The hot-dipped steel sheet of the present invention having excellentfatigue properties, elongation and collision properties, includes: thehigh-strength steel sheet of the present invention described above; anda hot-dipped layer provided on the surface of the high-strength steelsheet.

In the hot-dipped steel sheet of the present invention having excellentfatigue properties, elongation and collision properties, the hot-dippedlayer may consist of zinc.

The alloyed hot-dipped steel sheet of the present invention havingexcellent fatigue properties, elongation and collision properties,includes: the high-strength steel sheet of the present inventiondescribed above; and an alloyed hot-dipped layer provided on the surfaceof the high-strength steel sheet.

The method for producing the high-strength steel sheet of the presentinvention having excellent fatigue properties, elongation and collisionproperties, the method includes: heating a slab including: in terms ofpercent by mass %, 0.03 to 0.10% of C; 0.01 to 1.5% of Si; 1.0 to 2.5%of Mn; 0.1% or less of P; 0.02% or less of S; 0.01 to 1.2% of Al; 0.06to 0.15% of Ti; and 0.01% or less of N; and containing as the balance,iron and inevitable impurities, at a temperature in a range of 1,150 to1,280° C. and performing hot rolling under conditions where a finishrolling is finished at a temperature in a range of not less than an Ar₃point, thereby obtaining a hot-rolled material; coiling the hot-rolledmaterial in a temperature range of 600° C. or less, thereby obtaining ahot-rolled steel sheet; subjecting the hot-rolled steel sheet to acidpickling; subjecting the pickled hot-rolled steel sheet to first skinpass rolling at an elongation rate in a range of 0.1 to 5.0%; annealingthe hot-rolled steel sheet under conditions where a maximum heatingtemperature (Tmax° C.) is in a range of 600 to 750° C. and a holdingtime (t seconds) in a temperature range of 600° C. or higher fulfillsExpressions (1) and (2) as follows; and subjecting the annealedhot-rolled steel sheet to second skin pass rolling.

530−0.7×Tmax≦t≦3,600−3.9×Tmax  (1)

t>0  (2)

In the method for producing the high-strength steel sheet of the presentinvention having excellent fatigue properties, an elongation rate may beset to be in a range of 0.2 to 2.0% in the second skin pass rolling.

½ or more of the amount of Ti contained in the hot-rolled steel sheetafter the coiling may exist in a solid-solution state.

The method for producing the hot-dipped steel sheet of the presentinvention having excellent fatigue properties, elongation and collisionproperties, the method includes: heating a slab including: in terms ofpercent by mass %, 0.03 to 0.10% of C; 0.01 to 1.5% of Si; 1.0 to 2.5%of Mn; 0.1% or less of P; 0.02% or less of S; 0.01 to 1.2% of Al; 0.06to 0.15% of Ti; and 0.01% or less of N; and containing as the balance,iron and inevitable impurities, at a temperature in a range of 1,150 to1,280° C. and performing hot rolling under conditions where a finishrolling is finished at a temperature in a range of not less than an Ar₃point, thereby obtaining a hot-rolled material; coiling the hot-rolledmaterial in a temperature range of 600° C. or less, thereby obtaining ahot-rolled steel sheet; subjecting the hot-rolled steel sheet to acidpickling; subjecting the pickled hot-rolled steel sheet to first skinpass rolling at an elongation rate in a range of 0.1 to 5.0%; annealingthe hot-rolled steel sheet under conditions where a maximum heatingtemperature (Tmax° C.) is in a range of 600 to 750° C. and a holdingtime (t seconds) in a temperature range of 600° C. or higher fulfillsExpressions (1) and (2) as follows, and performing hot dipping to form ahot-dipped layer on a surface of the hot-rolled steel sheet, therebyobtaining a hot-dipped steel sheet; and subjecting the hot-dipped steelsheet to second skin pass rolling.

530−0.7×Tmax≦t≦3,600−3.9×Tmax  (1)

t>0  (2)

In the method for producing the hot-dipped steel sheet of the presentinvention having excellent fatigue properties, elongation and collisionproperties, an elongation rate may be set to be in a range of 0.2 to2.0% in the second skin pass rolling.

The method for producing the alloyed hot-dipped steel sheet of thepresent invention having excellent fatigue properties, elongation andcollision properties, the method includes: heating a slab comprising: interms of percent by mass %, 0.03 to 0.10% of C; 0.01 to 1.5% of Si; 1.0to 2.5% of Mn; 0.1% or less of P; 0.02% or less of S; 0.01 to 1.2% ofAl; 0.06 to 0.15% of Ti; and 0.01% or less of N; and containing as thebalance, iron and inevitable impurities, at a temperature in a range of1,150 to 1,280° C. and performing hot rolling under conditions where afinish rolling is finished at a temperature in a range of not less thanan Ar₃ point, thereby obtaining a hot-rolled material; coiling thehot-rolled material in a temperature range of 600° C. or less, therebyobtaining a hot-rolled steel sheet; subjecting the hot-rolled steelsheet to acid pickling; subjecting the pickled hot-rolled steel sheet tofirst skin pass rolling at an elongation rate in a range of 0.1 to 5.0%;annealing the hot-rolled steel sheet under conditions where a maximumheating temperature (Tmax° C.) is in a range of 600 to 750° C. and aholding time (t seconds) in a temperature range of 600° C. or higherfulfills Expressions (1) and (2) as follows, performing hot dipping toform a hot-dipped layer on a surface of the hot-rolled steel sheet so asto obtain a hot-dipped steel sheet, and subjecting the hot-dipped steelsheet to an alloying treatment to convert the hot-dipped layer into analloyed hot-dipped layer; and subjecting the hot-dipped steel sheet onwhich the alloying treatment is performed to second skin pass rolling.

530−0.7×Tmax≦t≦3,600−3.9×Tmax  (1)

t>0  (2)

In the method for producing the alloyed hot-dipped steel sheet of thepresent invention having excellent fatigue properties, elongation andcollision properties, an elongation rate may be set to be in a range of0.2 to 2.0% in the second skin pass rolling.

Effects of the Invention

In the method for producing the high-strength steel sheet of the presentinvention, a tensile strength in a range of 590 MPa or more is realizedby fulfilling the above-described component composition. In addition, Tiis added, and in the hot rolling stage, precipitation of alloycarbonitrides is suppressed by adjusting the coiling temperature, and inthe annealing stage, alloy carbonitrides are precipitated by adjustingthe heating temperature and the holding time. As a result, precipitationstrengthening is applied; and thereby, a high yield stress is realized.Therefore, a high collision energy absorbing ability (excellentcollision properties) can be achieved. In addition, by performing theskin pass before the annealing, strains are introduced only to thesurface layer of the steel sheet. This strains become precipitationsites of alloy carbonitrides during the annealing step; and therefore,precipitation of carbonitrides at or in the vicinity of the surfacelayer of the steel sheet can be accelerated during the annealing.Thereby, softening of the surface layer can be suppressed. As a result,Hvs/Hvc of the steel sheet can be set to be in a range of 0.85 or more;and thereby, high fatigue strength ratio (excellent fatigue properties)can be achieved. In addition, by performing the skin pass at apredetermined elongation rate, excellent elongation (excellentworkability) can be achieved.

Since the high-strength steel sheet of the present invention has theabove-described component composition and the microstructure, a tensilestrength in a range of 590 MPa or more and excellent elongation(excellent workability) can be realized. In addition, since a density ofTi(C,N) precipitates having sizes of 10 nm or smaller is in a range of10¹⁰ precipitates/mm³ or more, a high yield stress is realized.Therefore, a high collision energy absorbing ability (excellentcollision properties) can be achieved. In addition, since a ratio(Hvs/Hvc) is in a range of 0.85 or more, a high fatigue strength ratio(excellent fatigue properties) can be achieved.

The hot-dipped steel sheet of the present invention and the alloyedhot-dipped steel sheet of the present invention can achieve the sameeffects as those of the high-strength steel sheet described above andexcellent rust prevention.

Accordingly, the present invention can provide a high-strength steelsheet, a hot-dipped steel sheet, and an alloyed hot-dipped steel sheet,which have a tensile strength in a range of 590 MPa or more andexcellent fatigue properties, elongation and collision properties, andproduction methods thereof.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 is a graph showing a relationship between Hvs/Hvc and a fatiguestrength ratio.

FIG. 2 is a graph showing a relationship between an elongation rate offirst skin pass and Hvs/Hvc.

FIG. 3 is a graph showing a relationship between a tensile strength andan elongation.

FIG. 4 is a graph showing a relationship between a tensile strength anda fatigue strength ratio.

FIG. 5 is a graph showing a relationship between a maximum heatingtemperature (Tmax) of annealing and Hvs/Hvc.

FIG. 6 is a graph showing a relationship between a maximum heatingtemperature and a holding time in a temperature range of 600° C. orhigher during annealing.

FIG. 7 is a graph showing a relationship between an elongation rate(rolling rate) of a second skin pass after annealing and a fatiguestrength ratio.

FIG. 8 is a graph showing a relationship between Ti amount and ahardness ratio.

FIG. 9 is a graph showing a relationship between Ti amount and a yieldratio.

FIG. 10 is a graph showing a relationship between density of Ti(C,N)precipitates and a yield ratio.

FIG. 11 shows TEM photographs of the microstructure of ExperimentalExample B-k (steel of the present invention), FIG. 11( a) is aphotograph at 5,000-fold magnification, FIG. 11( b) is a photograph at100,000-fold magnification, and FIG. 11( c) is a photograph at100,000-fold magnification.

FIG. 12 shows TEM photographs of the microstructure of ExperimentalExample B-e (comparative steel), FIG. 12( a) is a photograph at5,000-fold magnification, and FIG. 12( b) is a photograph at500,000-fold magnification.

FIG. 13 is a graph showing a size distribution of Ti(C,N) ofExperimental Example B-k (steel of the present invention).

FIG. 14 is a graph showing a size distribution of Ti(C,N) ofExperimental Example B-e (comparative steel).

BEST MODE FOR CARRYING OUT THE INVENTION

Details of the present invention will be described below.

The inventors have focused on the fact that in order to produce ahigh-strength steel sheet, a hot-dipped steel sheet, or an alloyedhot-dipped steel sheet having excellent fatigue properties, elongation,and collision properties which cannot be achieved in the prior art,precipitation strengthening due to microalloying elements such as Ti,Nb, Mo, and V has to be utilized sufficiently, and have examinedinfluences of alloy components and production conditions onprecipitation behaviors.

That is, the inventors examined the precipitation behaviors of alloycarbonitrides of Ti, Nb, Mo, and V which occur during the production ofa high-strength steel sheet, a hot-dipped steel sheet, or an alloyedhot-dipped steel sheet. In detail, the inventors examined a coilingtemperature of a hot-rolled material, annealing conditions in anannealing step (including galvanization step), and an influence ofdislocations introduced to the surface of the steel sheet during skinpass rolling performed after acid-pickling the hot-rolled steel sheet.Then, the inventors examined an influence on fatigue properties,elongation, and collision properties.

As a result, the inventors found that in order to realize a high yieldstress by utilizing the precipitation strengthening for the purpose ofimproving collision properties, it is preferable to suppressprecipitation of alloy carbonitrides in a hot rolling stage and toprecipitate the alloy carbonitrides in a matrix so as to performprecipitation strengthening in an annealing stage. Further, theinventors thought that in order to increase the hardness of the surfacelayer of the steel sheet which has a large influence on the fatigueproperties, it is effective to precipitate the alloy carbonitrides at orin the vicinity of the surface layer of the steel sheet in the annealingstage. In addition, the inventors found that as a method foraccelerating precipitation of alloy carbonitrides, it is effective toperform skin pass rolling so as to intensively introduce strains only tothe surface layer and the vicinity thereof in the steel sheet afterperforming hot rolling and acid pickling. It is effective to increaseprecipitation sites of alloy carbonitrides by the skin pass rolling, andthese alloy carbonitrides precipitate during annealing; and thereby, anincrease in the strength is extended due to precipitation strengthening.In addition, the inventors also found that the surface roughness isimproved and the surface layer is work-hardened by subjecting the steelsheet to skin pass rolling at a rolling rate of 1.0% or more aftercompleting the annealing; and thereby, the fatigue properties arefurther improved.

Accordingly, it becomes possible to produce a steel sheet having a highyield stress which could not be achieved by a production method of ahigh-strength steel sheet, a hot-dipped steel sheet, or an alloyedhot-dipped steel sheet of the prior art. Specifically, by performingannealing after the skin pass rolling, the surface layer and thevicinity thereof are hardened by precipitation strengthening due to thealloy carbides; and thereby, fatigue properties are improved. Inaddition, by the skin pass rolling after the annealing, the surfaceroughness is further improved, and the surface layer and the vicinitythereof are work-hardened. Accordingly, the fatigue properties arefurther enhanced.

Next, the high-strength steel sheet of the present invention will bedescribed. At first, the reasons for limitations associated with thecomponents of the steel sheet are described.

The C content is set to be in a range of 0.03 to 0.10%. In the casewhere the C content is less than 0.03%, the strength is degraded, and590 MPa which is a target tensile strength cannot be achieved. Inaddition, a degree of hardening of the surface layer of the steel sheetafter annealing is reduced. Therefore, the C content is set to be in arange of 0.03% or more. On the other hand, in the case where the Ccontent exceeds 0.10%, the strength is increased excessively; andthereby, elongation is deteriorated drastically. Therefore, in practice,it becomes difficult to form, and furthermore, weldability isdeteriorated drastically. Therefore, the C content is set to be in arange of 0.10% or less.

The C content is preferably in a range of 0.06 to 0.09%. In this case, atensile strength of 590 MPa or more is obtained, and a fatigue strengthratio of 0.45 or more is also obtained.

Si is a solid-solution strengthening element and is effective inincreasing the strength; and therefore, as the Si content is increased,the balance between tensile strength and elongation is improved.However, when the Si content is too large, Si has an influence onwettability of galvanization and chemical conversion properties.Therefore, the upper limit of the Si content is set to be 1.5%. Inaddition, since Si is used for deoxidizing and Si is incorporatedinevitably, the lower limit thereof is set to be 0.01%.

It is preferable that the Si content be in a range of 1.2% or less.There may be cases where problems with wettability of galvanization orchemical conversion properties occur due to an influence of conditionsduring hot rolling or an atmosphere during continuous annealing.Therefore, the upper limit of the Si content is preferably 1.2%.

The Mn content is set to be in a range of 1.0 to 2.5%. Mn is aneffective element in enhancing solid-solution strengthening andhardenability; however, 590 MPa which is a target tensile strengthcannot be achieved in the case where the Mn content is less than 1.0%.Therefore, the Mn content is set to be in a range of 1.0% or more. Onthe other hand, in the case where the Mn content exceeds 2.5%,segregation is more likely to occur, and press formability isdeteriorated. In practice, the Mn content is preferably in a range of1.0 to 1.8% with regard to the steel sheet having a tensile strength of590 to 700 MPa, and the Mn content is preferably in a range of 1.6 to2.2% with regard to the steel sheet having a tensile strength of 700 MPato 900 MPa, and the Mn content is preferably in a range of 2.0 to 2.5%with regard to the steel sheet having a tensile strength of 900 MPa ormore. There is a suitable Mn amount range depending on the tensilestrength, and an excessive addition of Mn causes deterioration ofworkability due to Mn segregation. Therefore, it is preferable that theMn content be adjusted in accordance with the tensile strength asdescribed above.

P acts as a solid-solution strengthening element and increases thestrength of the steel sheet. However, when the P content is too large,workability or weldability of the steel sheet is degraded, which is notpreferable. In particular, in the case where the P content exceeds 0.1%,degradation of the workability or weldability of the steel sheet becomesnotable. Therefore, the P content is preferably set to be in a range of0.1% or less and is more preferably set to be in a range of 0.02% orless.

In the case where the S content is too large, inclusions such as MnS aregenerated; and thereby, stretch flangeability is degraded, andfurthermore, cracks occur during hot rolling. Therefore, it ispreferable that the S content be reduced to be as low as possible. Inparticular, in order to prevent the occurrence of cracks during hotrolling and obtain good workability, the S content is preferably set tobe in a range of 0.02% or less, and is more preferably set to be in arange of 0.01% or less.

The Al content is set to be in a range of 0.01 to 1.2%. By adding Al asa deoxidizing element, the amount of dissolved oxygen in a molten steelcan be efficiently reduced. In the case where the Al content is in arange of 0.01% or more, it is possible to prevent Ti, Nb, Mo, and Vwhich are important elements in the present invention from forming alloyoxides with dissolved oxygen. In this manner, Al is used fordeoxidizing; however, Al is incorporated inevitably. Therefore, thelower limit of the Al content is set to be 0.01%, and the Al content ispreferably in a range of 0.02% or more. On the other hand, in the casewhere the Al content exceeds 1.2%, Al becomes a factor that deterioratesgalvanizing properties and chemical conversion properties. Therefore,the Al content is set to be in a range of 1.2% or less and is preferablyset to be in a range of 0.6% or less.

Ti is an important element important in the present invention. Ti is animportant element for precipitation strengthening of the steel sheetduring annealing after hot rolling. In the production process, it isnecessary to maintain a solid solution state while suppressing theamount of formed precipitates as low as possible in a hot rolling stage(a stage from hot rolling to coiling); and therefore, a coilingtemperature during the hot rolling is set to be in a range of 600° C. orless at which Ti precipitates are less likely to be generated. Inaddition, skin pass rolling is performed before annealing; and thereby,dislocations are introduced. Next, in an annealing stage, Ti(C,N) isfinely precipitated on the introduced dislocations. In particular, at orin the vicinity of the surface layer of the steel sheet where adislocation density is increased, the effect (fine precipitation ofTi(C,N)) becomes notable. Due to this effect, it becomes possible toattain Hvs/Hvc 0.85, and high fatigue properties can be achieved. Inaddition, by precipitation strengthening due to an addition of Ti, ayield ratio which is a ratio between tensile strength and yield strengthcan be in a range of 0.80 or more. Among many precipitationstrengthening elements, Ti has the highest precipitation strengtheningability. This is because a difference between the solubility of Ti in aγ phase and the solubility of Ti in an α phase is large. In order toachieve a tensile strength of 590 MPa or more, Hvs/Hvc≧0.85, and a yieldratio of 0.80 or more, it is necessary to set the Ti content to be in arange of 0.06% or more as shown in FIGS. 8 and 9. In the case where theTi content is less than 0.06%, as shown in FIG. 10, a precipitatedensity of Ti(C,N) having sizes of 10 nm or smaller becomes less than10¹⁰ pieces/mm³; and thereby, a high yield ratio is not obtained. Ticontributes to precipitation strengthening, and in addition, Ti is anelement which delays a rate of recrystallization of austenite during hotrolling. Therefore, in the case where the Ti content is excessive, thetexture of the hot-rolled steel sheet is developed; and thereby,anisotropy after annealing is increased. In concrete, in the case wherethe Ti content exceeds 0.12%, the anisotropy of the steel sheet isincreased, and in the case where the Ti content exceeds 0.15%, theanisotropy of the steel sheet is particularly increased. As a result,workability is degraded. Therefore, the upper limit of the Ti content isset to be 0.15% and is preferably set to be 0.12%.

N forms TiN; and thereby, workability of the steel sheet is degraded.Therefore, it is preferable that the N content be as low as possible. Inparticular, in the case where the N content exceeds 0.01%, coarse TiN isgenerated; and thereby, the workability of the steel sheet isdeteriorated, and in addition, the amount of Ti which does notcontribute to precipitation strengthening is increased. Therefore, it ispreferable that the N content be set to be in a range of 0.01% or less.

The steel sheet of the present invention includes the above-describedelements and the balance which is iron and inevitable impurities. Asneeded, one or more selected from Nb, Mo, V, Ca, Mg, B, Cr, Cu, and Nidescribed as follows may further be contained.

Nb is an important element as a precipitation strengthening element likeTi. However, in the case where the Nb content is less than 0.005%, theeffect is small. Therefore, the lower limit of the Nb content is set tobe 0.005%. In addition, as is the case with Ti, Nb has an effect ofdelaying the rate of recrystallization of austenite during hot rolling.Therefore, in the case where the Nb content is excessive, workability isdeteriorated. In concrete, in the case where the Nb content exceeds0.1%, an increase in the strength by the precipitation strengthening issaturated, and in addition, elongation is degraded. Therefore, the upperlimit of the Nb content is set to be 0.1%. In the case where Nb iscontained together with Ti, the effect of making grain sizes finebecomes prominent. Therefore, it is preferable that the Nb content be ina range of 0.02 to 0.05%, and in this case, the above-described effectis obtained drastically.

As is the case with Ti and Nb, Mo and V are precipitation strengtheningelements. In the case where the Mo content and the V content are eachless than 0.005%, the effect is small. In addition, in the case wherethe Mo content and the V content each exceed 0.2%, the effect ofimproving the precipitation strengthening is small, and in addition,elongation is deteriorated. Therefore, the Mo content and the V contentare each set to be in a range of 0.005 to 0.2%.

Ca forms CaS which is a compound with S and is bonded to S. As a result,there is an effect of suppressing generation of MnS. Mg has an effect ofmaking inclusions fine. In the case where the Ca content and the Mgcontent each exceed 0.005%, the amount of inclusions is increased due tothe excessive addition; and thereby, hole expandability is deteriorated.Therefore, the upper limits thereof are set to be 0.005%. In addition,in the case where the Ca content and the Mg content are each less than0.0005%, the above-described effect is not sufficiently obtained.Therefore, it is preferable that the lower limits thereof be 0.0005%.

B is an element which can improve hardenability drastically. Therefore,in the case where sufficient cooling ability is not obtained due to thelimitation of equipment in a hot rolling line, or in the case wherecracks are generated in grain boundaries due to secondary workembrittlement, B is contained as needed for the purpose of strengtheninggrain boundaries. In the case where the B content exceeds 0.005%,improvement of the hardenability is not obtained in practice; andtherefore, the upper limit of the B content is set to be 0.005%. In thecase where the B content is less than 0.0005%, the above-describedeffect is not sufficiently obtained. Therefore, it is preferable thatthe lower limit of the B content be 0.0005%.

As is the case with Mn, Cr is one of elements effective in enhancinghardenability. Therefore, as the Cr content is increased, the tensilestrength of the steel sheet is increased. In the case where the Crcontent is large, Cr-based alloy carbides such as Cr₂₃C₆ areprecipitated, and when these carbides are preferentially precipitated inthe grain boundaries, press formability is deteriorated. Therefore, theupper limit of the Cr content is set to be 1%. In addition, in the casewhere the Cr content is less than 0.005%, the above-described effect isnot sufficiently obtained. Therefore, it is preferable that the lowerlimit of the Cr content be 0.005%.

Cu has an effect of increasing the strength of the steel material due toprecipitation thereof. Alloy elements such as Ti are bonded to C or Nand form alloy carbides; however, Cu is precipitated solely andstrengthens the steel material. However, a steel material containing alarge amount of Cu embrittles during hot rolling. Therefore, the upperlimit of the Cu content is set to be 1%. In addition, in the case wherethe Cu content is less than 0.005%, the above-described effect is notsufficiently obtained. Therefore, it is preferable that the lower limitof the Cu content be 0.005%.

As is the case with Mn, Ni enhances hardenability of the steel material,and in addition, Ni contributes to the improvement of toughness.Furthermore, Ni has an effect of preventing hot brittleness in the caseof including Cu. However, since alloy costs are very expensive, theupper limit of the Ni content is set to be 1%. In the case where the Nicontent is less than 0.005%, the above-described effect is notsufficiently obtained. Therefore, it is preferable that the lower limitof the Ni content be 0.005%.

Next, the microstructure of the steel sheet which is one of thecharacteristics of the present invention will be described.

According to the present invention, the microstructure includes bainiteat an area ratio of 40% or more and the balance being either one or bothof ferrite and martensite. Here, the microstructure is a microstructurein a sheet thickness center portion which is observed by taking a samplefrom a portion of the steel sheet that is ¼ of the sheet thickness innerfrom the surface.

In the present invention, in the case where the area ratio of bainite isin a range of 40% or more, an increase in the strength due toprecipitation strengthening can be expected. That is, a temperature atwhich the hot-rolled material is coiled is set to be in a range of 600°C. or less so as to ensure solid-solution Ti in the hot-rolled steelsheet, and this temperature is close to the bainite transformationtemperature. Therefore, a large amount of bainite is included in themicrostructure of the hot-rolled steel sheet, and transformationdislocations which area introduced simultaneously with transformationincrease an amount of TiC nucleation sites during annealing; andthereby, higher precipitation strengthening can be achieved. The arearatio of bainite is changed drastically due to a cooling history duringhot rolling; however, the area ratio of bainite is adjusted depending onthe needed material properties. The area ratio of bainite is preferablyin a range of more than 70%. In this case, the increase in the strengthdue to the precipitation strengthening is further enhanced, and inaddition, an amount of coarse cementite which is inferior in pressformability is reduced; and thereby, press formability can be maintainedproperly. The upper limit of the area ratio of bainite is preferably90%.

In the present invention, in the production process, in the hot rollingstage (a stage from hot rolling to coiling), Ti in the hot-rolled steelsheet is maintained in a solid-solution state, and then strains areintroduced to the surface layer by skin pass rolling after the hotrolling. Thereafter, in the annealing stage, Ti(C,N) is precipitated inthe surface layer while utilizing the introduced strains as nucleationsites. As a result, fatigue properties are improved. Therefore, it isimportant to complete (finish) the hot rolling in a temperature range of600° C. or less where precipitation of Ti is less likely to proceed.That is, it is important to coil the hot-rolled material at atemperature in a range of 600° C. or less. In the structure of thehot-rolled steel sheet obtained by coiling the hot-rolled material (thestructure in the hot rolling stage), the fraction of bainite may bearbitrary. In particular, in the case where high elongation is desiredfor products (high-strength steel sheet, hot-dipped steel sheet, andalloyed hot-dipped steel sheet), it is effective to increase thefraction of ferrite during hot rolling. On the other hand, in the casewhere hole expandability is considered to be important, the hot-rolledmaterial may be coiled at lower temperature; and thereby, themicrostructure including bainite and martensite as main phases may beformed.

As described above, since coiling is performed at a temperature in arange of 600° C. or less so as to ensure the amount of solid-solution Tiin the hot-rolled steel sheet, the microstructure of the hot-rolledsteel sheet (the microstructure in the hot rolling stage) substantiallyconsists of bainite and the balance being either one or both of ferriteand martensite. Thereafter, the hot-rolled steel sheet is heated to 600°C. or higher in the annealing; and thereby, bainite and martensite aretempered. In general, tempering means reducing a dislocation density bya heat treatment. Bainite and martensite generated at a temperature in arange of 600° C. or less are tempered during the annealing. Therefore,it can be said that bainite and martensite in the microstructure of theproducts are tempered bainite and tempered martensite in practice. Thetempered bainite and the tempered martensite are distinguished fromgeneral bainite and martensite because the tempered bainite and thetempered martensite have low dislocation densities as follows.

The microstructure of the hot-rolled steel sheet in the hot rollingstage contains bainite and martensite; and therefore, the dislocationdensity is high. However, since bainite and martensite are temperedduring the annealing, the dislocation density is reduced. In the casewhere an annealing time is insufficient, the dislocation density ismaintained at high value; and as a result, elongation becomes low.Therefore, it is preferable that the average dislocation density of thesteel sheet after annealing be in a range of 1×10¹⁴ m⁻² or less. In thecase where the annealing is performed under conditions that fulfillExpressions (1) and (2) described later, the reduction in thedislocation density proceeds simultaneously with precipitation ofTi(C,N). That is, in a state where precipitation of Ti(C,N) proceedssufficiently, the average dislocation density of the steel sheet isreduced. Typically, the reduction in the dislocation density causes areduction in the yield stress of the steel material. However, in thepresent invention, Ti(C,N) is precipitated simultaneously with thereduction in the dislocation density; and therefore, a high yield stressis obtained.

In the present invention, a measurement method of the dislocationdensity is performed on the basis of “a method of measuring adislocation density using X-ray diffraction” described in CAMP-ISIJ Vol.17 (2004) p. 396, and the average dislocation density is calculated fromthe half-value widths of diffraction peaks of (110), (211), and (220).

Since the microstructure has the above-described properties, a highyield ratio and a high fatigue strength ratio can be achieved which arenot achieved by a steel sheet that is produced by utilizingprecipitation strengthening in the prior art. That is, even in the casewhere the microstructure at or in the vicinity of the surface layer ofthe steel sheet includes ferrite as a main phase and exhibits a coarsestructure unlike the microstructure in the sheet thickness centerportion, the hardness of the surface layer and the vicinity thereof inthe steel sheet reaches a hardness substantially equivalent to that ofthe center portion of the steel sheet due to the precipitation ofTi(C,N) during annealing. As a result, generation of fatigue cracks issuppressed; and thereby, the fatigue strength ratio is increased.

Next, the reason for limitations associated with the tensile strength ofthe steel sheet which is the feature of the present invention will bedescribed.

The tensile strength of the steel sheet of the present invention is in arange of 590 MPa or more. The upper limit of the tensile strength is notparticularly limited. However, in a component range of the presentinvention, the upper limit of the practical tensile strength is about1180 MPa.

Here, the tensile strength is evaluated by the following method. A No. 5specimen described in JIS-Z2201 is produced, and then a tensile test isperformed according to a test method described in JIS-Z2241.

In the present invention, a ratio (yield ratio) of the yield strength tothe tensile strength which are obtained by the tensile test becomes 0.80or more due to precipitation strengthening.

In order to attain a high yield ratio as in the present invention,precipitation strengthening due to Ti(C,N) and the like which isprecipitated by the tempering of bainite is more important thantransformation strengthening due to a hard phase such as martensite. Inthe present invention, a density of Ti(C,N) precipitates having sizes of10 nm or smaller which is effective in precipitation strengthening is ina range of 10¹⁰ pieces/mm³ or more. Thereby, a yield ratio in a range of0.80 or more described above can be realized. Here, precipitates ofwhich the equivalent circular diameter obtained by a square root of(major axis×minor axis) is larger than 10 nm does not have an influenceon the properties obtained in the present invention. In contrast, as thesize of the precipitate becomes smaller, precipitation strengthening dueto Ti(C,N) is obtained more effectively; and as a result, there is apossibility that an added amount of alloy elements can be reduced.Therefore, a density of Ti(C,N) precipitates having grain sizes of 10 nmor smaller is defined.

Here, the precipitates are observed by the following method. A replicasample is produced according to a method described in Japanese PatentApplication, First Publication No. 2004-317203, and then the replicasample is observed with a transmission electron microscope. Themagnification of the field of view is set to be in a range of 5,000-foldmagnification to 100,000-fold magnification, and the number of Ti(C,N)having sizes of 10 nm or smaller is counted from 3 or more fields ofview. In addition, an electrolytic weight is obtained from a change inweight before and after electrolysis, and the weight is converted into avolume by a specific gravity of 7.8 ton/m³. Then, the counted number isdivided by the volume; and thereby, the precipitation density iscalculated.

Next, the reasons for limitations associated with a hardnessdistribution of the steel sheet which is one of the characteristics ofthe present invention will be described.

The inventors have found that in order to improve fatigue properties,elongation, and collision properties in a high-strength steel sheet inwhich precipitation strengthening due to microalloying elements isutilized, fatigue properties are improved by setting a ratio of thehardness of the surface layer of the steel sheet to the hardness of thecenter portion of the steel sheet to be in a range of 0.85 or more.Here, the hardness of the surface layer of the steel sheet is a hardnessat a portion that is 20 μm (at a depth of 20 μm) inner from the surfaceand is represented by Hvs. In addition, the hardness of the centerportion of the steel sheet is a hardness at a portion that is ¼ of thesheet thickness (at a depth of ¼ of the sheet thickness) inner from thesurface of the steel sheet and is represented by Hvc. The inventors havefound that the fatigue properties are deteriorated in the case where theratio Hvs/Hvc is less than 0.85, and on the other hand, the fatigueproperties are improved in the case where the ratio Hvs/Hvc is 0.85 ormore. Therefore, Hvs/Hvc is set to be in a range of 0.85 or more.

FIG. 1 shows a relationship between Hvs/Hvc and fatigue strength ratio.It can be seen that a fatigue strength ratio of 0.45 or more can beachieved in the case where Hvs/Hvc is in a range of 0.85 or more.Therefore, high fatigue properties are obtained. Here, in the case ofthe hot-dipped steel sheet or the alloyed hot-dipped steel sheet, thesurface layer means a range excluding the plating thickness. That is,the hardness of the surface layer is a hardness at a portion which isnot included in a hot-dipped layer or an alloyed hot-dipped layer andwhich is 20 μm inner from the surface of the high-strength steel sheet.In addition, the reason of determining the measurement portion of thehardness of the surface layer of the steel sheet to a portion that is 20μm (at a depth of 20 μm) inner from the surface is described as follows.In practice, with regard to a steel sheet having a tensile strength of590 MPa or more, the hardness is measured in a cross-section of thesteel sheet using a Vickers hardness tester. Based on the premise ofthis measurement, the measurement portion is determined from themeasurement ability. Therefore, in the case where it is possible tomeasure the hardness of the surface layer at a portion further closer tothe surface by using a nanoindentation technique, the measurementportion may be determined based on the measurement ability. Here, in thecase where measurement is performed at a portion different from theportion that is 20 μm (at a depth of 20 μm) inner from the surface, itis impossible to simply compare the absolute values of the measured Hvsand Hvc since the measurement methods are different. However, thethreshold of Hvs/Hvc which is a ratio of these harnesses can be used asit is.

In the present invention, the type of the steel sheet which is a productis a high-strength steel which is obtained by subjecting a hot-rolledsteel sheet to acid pickling and skin pass rolling and thereafterperforming annealing thereon.

The hot-dipped steel sheet of the present invention includes theabove-described high-strength steel sheet of the present invention, andthe hot-dipped layer provided on the surface of the high-strength steelsheet. In addition, the alloyed hot-dipped steel sheet of the presentinvention includes the above-described high-strength steel sheet of thepresent invention, and the alloyed hot-dipped layer provided on thesurface of the high-strength steel sheet.

As the hot-dipped layer and the alloyed hot-dipped layer, for example,layers consisting of either one or both of zinc and aluminum may beemployed, and specifically, a hot-dip galvanized layer, an alloyedhot-dip galvanized layer, a hot-dip aluminized layer, an alloyed hot-dipaluminized layer, a hot-dip Zn—Al coated layer, an alloyed hot-dip Zn—Alcoated layer, and the like may be employed. In particular, in terms ofplatability and corrosion resistance, a hot-dip galvanized layer and analloyed hot-dip galvanized layer which consist of zinc are preferable.

The hot-dipped steel sheet or the alloyed hot-dipped steel sheet areproduced by subjecting the above-described high-strength steel sheet ofthe present invention to hot dipping or alloyed hot-dipping. Here, thealloyed hot-dipping is a process of performing hot dipping to produce ahot-dipped layer on the surface and performing an alloying treatmentthereon to make the hot-dipped layer into an alloyed hot-dipped layer.

The hot-dipped steel sheet or the alloyed hot-dipped steel sheetincludes the high-strength steel sheet of the present invention, and thehot-dipped layer or the alloyed hot-dipped layer is formed on thesurface; and therefore, the effects of the high-strength steel sheet ofthe present invention and excellent rust prevention can be achieved.

Next, a method for manufacturing the high-strength steel sheet of thepresent invention will be described.

First, a slab having the above-described component composition isre-heated at a temperature in a range of 1,150 to 1,280° C. As the slab,a slab immediately after being produced by continuous casting equipment,or a slab produced by an electric furnace may be used.

By setting the heating temperature of the slab to be in a range of1,150° C. or more, carbide-forming elements and carbon can besufficiently decomposed and dissolved into the steel material. However,in the case where the heating temperature of the slab exceeds 1,280° C.,it is not preferable in terms of production costs; and therefore, theupper limit is set to be 1,280° C. In order to dissolve precipitatedcarbonitrides, it is preferable that the heating temperature be in arange of 1,200° C. or more.

Next, the re-heated slab is subjected to hot rolling under conditionswhere finish rolling is finished at a temperature in a range of the Ar₃point or more; and thereby, a hot-rolled material is obtained. Then, thehot-rolled material is coiled in a temperature range of 600° C. or less;and thereby, a hot-rolled steel sheet is obtained.

In the case where a finishing temperature (a temperature at which finishrolling is finished) during the hot rolling is less than the Ar₃ point,precipitation of alloy carbonitrides or coarsening of grains proceeds inthe surface layer; and thereby, the strength of the surface layerreduces notably. Therefore, excellent fatigue properties are notobtained. Consequently, in order to prevent deterioration of the fatigueproperties, the lower limit of the finishing temperature during the hotrolling is set to be in a range of Ar₃ point or more. The upper limit ofthe finishing temperature is not particularly limited; however, inpractice, the upper limit thereof is about 1,050° C.

Next, a cooling history from the finishing temperature during the hotrolling to the coiling will be described.

In the present invention, by setting the coiling temperature to be in arange of 600° C. or less, precipitation of alloy carbonitrides in thestage of the hot-rolled steel sheet (the stage from hot rolling tocoiling) is suppressed. The coiling temperature is important, and theproperties of the present invention are not degraded by the coolinghistory before the start of the coiling.

However, in the case where the ratio of the microstructure is adjustedso as to set the balance between elongation and hole expandability,which are mainly used as indexes of formability of a steel sheet for anautomobile, to a desired value, it is necessary to control the coolinghistory from the finishing temperature to the start of coiling. Forexample, as a fraction of ferrite is increased, elongation is improved;however, hole expandability is deteriorated.

Therefore, in the case where a steel sheet is produced of whichelongation is considered to be important, it is necessary to reduce thefinishing temperature and to conduct air cooling in a temperature rangeimmediately above a bainite starting temperature (Bs point) so as tocause ferrite transformation positively. In particular, it is preferableto positively cause ferrite transformation during hot rolling.Specifically, the finishing temperature is set to be in a range of theAr₃ point or more to (Ar₃ point+50° C.) or less; and thereby, a lot ofprocessing strains are introduced to austenite before transformation.Then, these strains are utilized as nucleation sites of ferrite, and atemperature is held in a temperature range in which ferritetransformation is most likely to proceed, specifically, from 600 to 680°C. for 1 to 10 seconds. In this manner, it is preferable that ferritetransformation be accelerated. After this intermediate holding, it isnecessary to cool again and to coil in a temperature range of 600° C. orless.

On the other hand, in the case where a steel sheet is produced of whichhole expandability is considered to be important, it is effective toincrease the finishing temperature and to perform rapid cooling to atemperature in a range of the Bs point or less in order to increasehardenability. In particular, it is preferable that the microstructurebe more homogeneous and mechanical properties thereof have lessanisotropy. Specifically, the finishing temperature is set to be in arange of (Ar₃+50° C.) or more; and thereby, the orientation of crystalsis arranged with a specific direction during hot rolling. As a result,the development of texture is suppressed. In addition, it is preferablethat in order to form a bainite single-phase structure, the coilingtemperature of the hot-rolled material be in a range of 300 to 550° C.

In the case where the coiling temperature exceeds 600° C., precipitationof alloy carbonitrides proceeds in the hot-rolled steel sheet.Therefore, the increase in the strength due to precipitate strengtheningafter annealing is not sufficiently obtained, and fatigue properties aredeteriorated. Accordingly, the upper limit of the coiling temperature isset to be 600° C. The lower limit is not particularly provided. As thecoiling temperature is lowered, amounts of solid-solubilized Ti, Nb, Mo,and V are increased; and thereby, the increase in the strength due toprecipitation strengthening during annealing is enhanced. Therefore, inorder to obtain the properties of the present invention, a lower coilingtemperature is effective. However, in practice, since the steel sheet iscooled by water cooling, the room temperature becomes the lower limit.

As described above, during the hot rolling stage, the coilingtemperature is controlled so as to suppress precipitation of alloycarbonitrides; and thereby, Ti maintains in a solid-solution state whilesuppressing the amount of formed precipitates as low as possible. In thehot-rolled steel sheet after coiling, it is preferable that ½ or greaterof the amount of contained Ti exists in the solid-solution state. Inthis case, the increase in the strength due to precipitationstrengthening after annealing is further enhanced.

Next, the hot-rolled steel sheet is pickled, and then the pickledhot-rolled steel sheet is subjected to first skin pass rolling at anelongation rate in a range of 0.1 to 5.0%.

The reason for limitations of the elongation during the first skin passrolling after acid pickling is described.

In the present invention, it is an important production condition toperform the first skin pass at an elongation in a range of 0.1 to 5.0%.By subjecting the hot-rolled steel sheet to skin pass, strains areprovided in the surface of the steel sheet. During annealing in asubsequent step, nuclei of alloy carbonitrides are more likely to beformed on the dislocation via these strains; and thereby, the surfacelayer is hardened. In the case where the elongation rate of the skinpass is less than 0.1%, sufficient strains cannot be provided; and as aresult, the surface layer hardness Hvs is not increased. On the otherhand, in the case where the elongation rate of the skin pass exceeds5.0%, strains are provided not only in the surface layer but also in thecenter portion of the steel sheet; and as a result, the workability ofthe steel sheet is degraded. In a typical steel sheet, ferrite isrecrystallized by the subsequent annealing; and thereby, elongation orhole expandability is improved. However, in the case where the componentcomposition of the present invention is included and coiling isperformed in a temperature range of 600° C. or less, Ti, Nb, Mo, and Vwhich are solid-solubilized in the hot-rolled steel sheet drasticallydelay ferrite recrystallization due to annealing; and thereby,elongation and hole expandability after annealing is not improved.Therefore, the upper limit of the elongation rate of the skin passrolling is set to be 5.0%. Strains are provided in accordance with theelongation rate of the skin pass rolling. In terms of improvement offatigue properties, precipitation strengthening proceeds in the surfacelayer and the vicinity thereof in the steel sheet during annealing inaccordance with the amount of strains in the surface layer of the steelsheet. Therefore, it is preferable that the elongation rate be in arange of 0.4% or more. In addition, in terms of workability of the steelsheet, in order to prevent deterioration of the workability due to thestrains provided in the steel sheet, it is preferable that theelongation rate be in a range of 2.0% or less.

From the results of FIG. 2, it can be identified that in the case wherethe elongation rate of the skin pass rolling is in a range of 0.1 to5.0%, Hvs/Hvc is improved to be in a range of 0.85 or more. In addition,it can also be identified that in the case where skin pass is notperformed (the elongation rate of the skin pass rolling is 0%), or inthe case where the elongation rate of the skin pass rolling exceeds 5%,Hvs/Hvc<0.85 is fulfilled.

From the results of FIG. 3, it can be identified that in the case wherethe elongation rate of the first skin pass is in a range of 0.1 to 5.0%,excellent elongation is obtained. In addition, it can also be identifiedthat in the case where the first skin pass elongation rate exceeds 5.0%,elongation is deteriorated, and press formability is deteriorated. Fromthe results of FIG. 4, it can be identified that in the case where thefirst skin pass rate is 0% or exceeds 5%, the fatigue strength ratio isdeteriorated.

From the results of FIGS. 3 and 4, it can be identified that in the casewhere the elongation rate of the skin pass rolling is in a range of 0.1to 5.0%, substantially the same elongation and fatigue strength ratioare obtained if tensile strengths are substantially the same. It can beidentified that in the case where the elongation rate of the skin passrolling exceeds 5% (high skin pass region), elongation is low and thefatigue strength ratio is also low, compared to those of the steel sheetof the present invention having a tensile strength in the same level.

Next, the hot-rolled steel sheet is annealed after performing the firstskin pass rolling. In addition, for the purpose of shape correction,leveling may be used.

In the present invention, the purpose of performing annealing is not totemper the hard phase but to precipitate Ti, Nb, Mo, and V as alloycarbonitrides from Ti, Nb, Mo, and V which are solid-solubilized(dissolved as a solid solution) in the hot-rolled steel sheet.Accordingly, it is important to control a maximum heating temperature(Tmax) and a holding time during the annealing step. The maximum heatingtemperature and the holding time are controlled to be in predeterminedranges; and thereby, not only the tensile strength and the yield stressare increased, but also the surface layer hardness is enhanced. As aresult, the fatigue properties and collision properties are improved. Inthe case where the temperature and the holding time during annealing areinappropriate, carbonitrides are not precipitated or precipitatedcarbonitrides coarsen. Therefore, the maximum heating temperature andthe holding time are limited as follows.

In the present invention, the maximum heating temperature duringannealing is set to be in a range of 600 to 750° C. In the case wherethe maximum heating temperature is less than 600° C., a time required toprecipitate alloy carbonitrides becomes long drastically; and thereby,it becomes difficult to produce the steel sheet in continuous annealingequipment. Therefore, the lower limit thereof is set to be 600° C. Inaddition, in the case where the maximum heating temperature exceeds 750°C., coarsening of alloy carbonitrides occurs; and thereby, the increasein the strength due to precipitation strengthening is not sufficientlyobtained. In addition, in the case where the maximum heating temperatureis in a range of an Ac₁ point or more, the temperature is in a two-phaseregion of ferrite and austenite; and thereby, the increase in strengthdue to the precipitate strengthening is not sufficiently obtained.Therefore, the upper limit thereof is set to be 750° C. The main purposeof the annealing is not to temper the hard phase but to precipitate Tiwhich is solid-solubilized in the hot-rolled steel sheet. Here, thefinal strength is determined by alloy components of the steel materialand the fraction of each phase in the microstructure of the hot-rolledsteel sheet. However, the improvement of the fatigue properties due tothe hardening of the surface layer and the enhancement of the yieldratio, which are the characteristics of the present invention, are notinfluenced by the alloy components of the steel material and thefraction of each phase in the microstructure of the hot-rolled steelsheet.

As a result of the tests, it was found that in the case where a holdingtime (t) in a temperature range of 600° C. or higher during annealingfulfills a relationship of Expressions (1) and (2) as follows inrelation to the maximum heating temperature Tmax during annealing, ahigh yield stress and Hvs/Hvc in a range of 0.85 or more are attained.

530−0.7×Tmax≦t≦3,600−3.9×Tmax  (1)

t>0  (2)

From the results of FIG. 5, it can be identified that in the case wherethe maximum heating temperature is in a range of 600 to 750° C., Hvs/Hvcbecomes 0.85 or more.

Moreover, as shown in FIG. 6, all the steel sheets of the presentinvention in examples are produced under conditions where the holdingtime (t) in a temperature range of 600° C. or higher fulfills the rangesof the Expressions (1) and (2). From the evaluation results of the steelsheets of the present invention in the examples, it can be identifiedthat in the case where the holding time (t) fulfills the ranges ofExpressions (1) and (2), Hvs/Hvc becomes 0.85 or more.

From the examples, it can be identified that in the case where Hvs/Hvcis in a range of 0.85 or more, the fatigue strength ratio becomes 0.45or more. In the case where the maximum heating temperature is in a rangeof 600 to 750° C., the surface layer is hardened due to precipitationstrengthening; and thereby, Hvs/Hvc becomes 0.85 or more. By setting themaximum heating temperature and the holding time in a temperature rangeof 600° C. or higher to be in the above-described ranges, the surfacelayer is sufficiently hardened compared to the hardness of the centerportion of the steel sheet. As a result, as shown in the examples, thefatigue strength ratio becomes 0.45 or more. This is because generationof fatigue cracks can be delayed by the hardening of the surface layer.As the surface layer hardness is increased, the effect is increased.

In addition, from the results of FIG. 5, it can be identified that inthe case where the maximum heating temperature is not in the range (outof the range) of 600 to 750° C., Hvs/Hvc<0.85 is fulfilled. In addition,from the examples, it can be identified that even in the case where themaximum heating temperature is in a range of 600 to 750° C.,Hvs/Hvc<0.85 is fulfilled if the coiling temperature of the hot-rolledmaterial and the elongation rate of the skin pass are not in the rangesof the present invention.

Thereafter, the annealed hot-rolled steel sheet is subjected to secondskin pass rolling. Thereby, the fatigue properties can further beimproved.

During the second skin pass rolling, the elongation rate is preferablyset to be in a range of 0.2 to 2.0%, and the elongation rate is morepreferably in a range of 0.5 to 1.0%. In the case where the elongationrate is less than 0.2%, a surface roughness is not improved sufficientlyand work hardening of only the surface layer is not proceeded. As aresult, there may be cases where fatigue properties are not sufficientlyimproved. Therefore, it is preferable that the lower limit thereof isset to be 0.2%. On the other hand, in the case where the elongation rateexceeds 2.0%, the steel sheet is hardened too much; and as a result,there may be cases where press formability is deteriorated. In addition,for example, among examples described later, in Experimental ExampleL-a, since the elongation rate of the second skin pass rolling afterannealing is 2.5%, the elongation becomes 17% which is inferior to thoseof other Experimental Examples. There may be cases where the elongationis degraded as is the case with Experimental Example L-a. Therefore, itis preferable that the upper limit be 2.0%.

The component composition containing alloying elements and productionconditions are controlled precisely in the above-described manner; andthereby, a high-strength steel sheet can be produced which has excellentfatigue properties and collision safety that cannot be achieved in theprior art and has a tensile strength in a range of 590 MPa or more.

The method for manufacturing the hot-dipped steel sheet of the presentinvention includes: a step of producing a hot-rolled steel sheet as isthe case with the above-described method for manufacturing thehigh-strength steel sheet of the present invention; a step ofacid-pickling the hot-rolled steel sheet; a step of subjecting thehot-rolled steel sheet to first skin pass rolling at an elongation ratein a range of 0.1 to 5.0%; a step of annealing the hot-rolled steelsheet under conditions where a maximum heating temperature (Tmax° C.) isin a range of 600 to 750° C. and a holding time (t seconds) in atemperature range of 600° C. or higher fulfills the Expressions (1) and(2), and performing hot dipping to form a hot-dipped layer on a surfaceof the hot-rolled steel sheet, thereby obtaining a hot-dipped steelsheet; and a step of subjecting the hot-dipped steel sheet to secondskin pass rolling.

The step until the hot-rolled steel sheet is obtained, the step ofacid-pickling, the step of performing the first skin pass rolling, andthe annealing are performed under the same conditions as those of theabove-described method for manufacturing the high-strength steel sheetof the present invention.

The conditions of the hot dipping are not particularly limited, and awell-known technique is applied. As a kind of plating elements, forexample, either one or both of zinc and aluminum may be employed.

During the second skin pass rolling, the elongation rate is preferablyset to be in a range of 0.2 to 2.0%, and the elongation rate is morepreferably in a range of 0.5 to 1.0%. Thereby, as shown in FIG. 7, thefatigue strength is further improved, and the fatigue strength ratio canfurther be improved. It is thought that this is because the surfacelayer is further hardened by the work hardening of the surface layer ofthe steel sheet due to the skin pass rolling. In the case where theelongation rate is less than 0.2%, there may be cases where sufficientwork hardening is not obtained. Therefore, it is preferable that thelower limit thereof is set to be 0.2%. In the case where the elongationrate exceeds 2.0%, there may be cases where the improvement of thefatigue strength ratio is not confirmed, and furthermore, there may alsobe cases where the elongation is degraded. Therefore, it is preferablethat the lower limit be 2.0%.

The method for manufacturing an alloyed hot-dipped steel sheet of thepresent invention includes: a step of producing a hot-rolled steel sheetas is the case with the above-described method for manufacturing thehigh-strength steel sheet of the present invention; a step ofacid-pickling the hot-rolled steel sheet; a step of subjecting thehot-rolled steel sheet to first skin pass rolling at an elongation ratein a range of 0.1 to 5.0%; a step of annealing the hot-rolled steelsheet under conditions where a maximum heating temperature (Tmax° C.) isin a range of 600 to 750° C. and a holding time (t seconds) in atemperature range of 600° C. or higher fulfills the Expressions (1) and(2), performing hot dipping to form a hot-dipped layer on a surface ofthe hot-rolled steel sheet, thereby obtaining a hot-dipped steel sheet,and subjecting the hot-dipped steel sheet to an alloying treatment toconvert the hot-dipped layer into an alloyed hot-dipped layer; and astep of subjecting the hot-dipped steel sheet on which the alloyingtreatment is performed to second skin pass rolling.

The step until the hot-rolled steel sheet is obtained, the step ofacid-pickling, the step of performing the first skin pass rolling, andthe annealing are performed under the same conditions as those of theabove-described method for manufacturing the high-strength steel sheetof the present invention. In addition, the step of performing hotdipping is performed under the same conditions as those of theabove-described method for manufacturing the hot-dipped steel sheet ofthe present invention.

The conditions of the alloying treatment are not particularly limited,and a well-known technique is applied.

During the second skin pass rolling, the elongation rate is preferablyset to be in a range of 0.2 to 2.0%, and the elongation rate is morepreferably in a range of 0.5 to 1.0%. Thereby, the fatigue strengthratio can further be improved. In the case where the elongation rate isless than 0.2%, there may be cases where sufficient work hardening isnot obtained. Therefore, it is preferable that the lower limit thereofis 0.2%. In the case where the elongation rate exceeds 2.0%, there maybe cases where the improvement of the fatigue strength ratio is notconfirmed, and furthermore, there may also be cases where the elongationis degraded. Therefore, it is preferable that the lower limit be 2.0%.

EXAMPLES

Hereinafter, examples of the present invention are described.

Using steel materials (slabs) Nos. A to Z shown in Table 1, steel sheetswere produced under the condition shown in Tables 2 to 8. Here, Ar₃ inTable 1 is a value calculated by Expression (3) as follows. Thecompositional ratios (the content of each element) are all representedby mass %, and underlined values represent out of the range of thepresent invention.

Ar₃=910−310×C−80×Mn−80×Mo+33×Si+40×Al  (3)

Here, element symbols in Expression (3) represent the contents (mass %)of the elements.

TABLE 1 Steel No. C Si Mn P S Al N Ti Nb Mo V Ca Mg B Ar3 Note A 0.040.04 1.34 0.0103 0.0045 0.04 0.0036 0.069 — — — — — — 791 Steel ofInvention B 0.06 0.18 1.95 0.0076 0.0040 0.03 0.0044 0.085 0.030 — — — —— 731 Steel of Invention C 0.08 0.65 2.30 0.0082 0.0035 0.03 0.00380.135 0.025 — — — — — 681 Steel of Invention D 0.06 0.52 2.06 0.00960.0062 0.03 0.0051 0.112 0.040 — 0.005 — 0.0016 — 711 Steel of InventionE 0.09 1.00 2.05 0.0085 0.0039 0.03 0.0035 0.065 — 0.150 — — — — 674Steel of Invention F 0.05 0.03 1.65 0.0095 0.0042 0.62 0.0038 0.068 — —0.030 — — 0.0012 786 Steel of Invention G 0.07 0.52 1.68 0.0085 0.00550.03 0.0034 0.078 0.044 — — 0.0013 — — 738 Steel of Invention H 0.080.46 1.23 0.0073 0.0067 0.04 0.0035 0.063 — — — — — — 773 Steel ofInvention I 0.07 0.13 1.85 0.0055 0.0035 0.03 0.0045 0.072 0.090 — — — —— 737 Steel of Invention J 0.06 0.18 1.75 0.0082 0.0044 0.04 0.00350.092 0.075 — — — — 0.0015 747 Steel of Invention K 0.07 0.15 2.010.0079 0.0066 0.04 0.0035 0.102 0.036 0.003 — 0.0015 — — 724 Steel ofInvention L 0.08 1.06 2.45 0.0085 0.0056 0.02 0.0038 0.142 0.031 — 0.0030.0011 — 0.0013 655 Steel of Invention M 0.02 0.02 1.81 0.0081 0.00340.03 0.0042 0.065 — — — — — — 761 Comparative Steel N 0.15 0.53 2.300.0091 0.0035 0.02 0.0049 0.080 — — — 0.0010 — — 698 Comparative Steel O0.06 1.65 1.25 0.0053 0.0041 0.03 0.0034 0.075 0.021 0.003 0.012 — — —847 Comparative Steel P 0.08 0.03 0.72 0.0054 0.0045 0.03 0.0029 0.0720.053 — 0.051 — — — 830 Comparative Steel Q 0.06 0.03 2.70 0.0068 0.00380.02 0.0038 0.065 0.041 0.032 0.058 — 0.0022 — 675 Comparative Steel R0.09 0.04 0.95 0.0081 0.0052 1.72 0.0039 0.075 0.051 0.021 0.064 — — —875 Comparative Steel S 0.06 0.15 1.68 0.0102 0.0053 0.30 0.0034 0.042 —— — — — — 773 Comparative Steel T 0.09 0.52 2.44 0.0072 0.0059 0.140.0051 0.186 — — 0.002 — — 0.0016 725 Comparative Steel

hot rolling, coiling, acid pickling, first skin pass rolling, annealing,and second skin pass were performed in this order; and thereby,high-strength steel sheets were produced. All the sheet thicknesses ofhot-rolled materials after the hot rolling were set to be 3.0 mm. Therate of temperature increase during the annealing was set to be 5° C./s,and the rate of cooling from the maximum heating temperature was set tobe 5° C./s.

In addition, for several Experimental Examples, galvanization and analloying treatment were performed after the annealing to produce hot-dipgalvanized steel sheets and alloyed hot-dip galvanized steel sheets.Here, in the case where the hot-dip galvanized steel sheets wereproduced, second skin pass was performed after the hot-dipgalvanization, and in the case where the alloyed hot-dip galvanizedsteel sheets were produced, second skin pass was performed after thealloying treatment.

TABLE 2 First skin Annealing Hot rolling pass Maximum Heating FinishingCooling Coiling Elongation heating Holding Left side of Right side ofExperimental Steel temperature temperature rate temperature ratetemperature time Expression Expression Example No. (° C.) (° C.) (°C./s) (° C.) (%) (° C.) (sec) (1) (° C.) (1) (° C.) A-a A 1230 910 25515 0.8 650 240 75 1065 A-b 1235 915 50 510 1.5 720 120 26 792 B-a B1220 905 45 520 0.5 680 240 54 948 B-b 1220 920 45 530 0.5 700 60 40 870C-a C 1220 895 40 510 0.5 690 240 47 909 C-b 1220 890 40 425 0.3 700 8040 870 D-a D 1225 900 35 520 0.5 660 120 68 1026 D-b 1220 895 35 525 0.5680 320 54 948 E-a E 1210 905 50 515 0.5 660 300 68 1026 E-b 1210 910 50530 0.5 660 95 68 1026 F-a F 1220 895 40 525 0.5 660 300 68 1026 F-b1220 895 45 510 0.5 670 75 61 987 G-a G 1230 920 45 500 0.5 680 120 54948 G-b 1225 920 20 530 1.5 720 200 26 792 H-a H 1220 920 45 520 0.8 630480 89 1143 H-b 1200 880 40 530 2.5 680 260 54 948 I-a I 1220 930 45 5100.8 700 240 40 870 I-b 1225 920 50 520 0.5 710 120 33 831 J-a J 1225 89045 480 0.8 710 680 33 831 J-b 1220 910 45 480 0.8 650 240 75 1065

TABLE 3 First skin Annealing Hot rolling pass Maximum Heating FinishingCooling Coiling Elongation heating Holding Left side of Right side ofExperimental Steel temperature temperature rate temperature ratetemperature time Expression Expression Example No. (° C.) (° C.) (°C./s) (° C.) (%) (° C.) (sec) (1) (° C.) (1) (° C.) K-a K 1200 900 50500 0.8 690 80 47 909 K-b 1230 910 35 450 0.8 680 600 54 948 L-a L 1220920 40 550 0.5 710 180 33 831 L-b 1225 890 45 500 0.8 690 600 47 909 M-aM 1215 900 40 510 0.8 650 120 75 1065 M-b 1210 910 45 520 0.8 680 120 54948 N-a N 1205 910 40 140 0.5 680 400 54 948 N-b 1200 920 40 510 0.8 680890 54 948 O-a O 1210 905 45 450 0.5 680 100 54 948 O-b 1210 915 45 5000.5 700 600 40 870 P-a P 1230 915 45 450 0.5 680 240 54 948 P-b 1230 91545 480 0.5 650 600 75 1065 Q-a Q 1210 890 50 480 0.8 710 200 33 831 Q-b1210 895 40 490 0.8 700 260 40 870 R-a R 1225 905 40 550 0.5 650 200 751065 R-b 1225 920 45 500 0.5 680 200 54 948 S-a S 1210 910 40 550 0.4670 240 61 987 S-b 1210 905 40 520 0.4 670 120 61 987 T-a T 1220 910 40480 0.5 710 240 33 831 T-b 1220 910 50 490 0.6 700 200 40 870

TABLE 4 Experi- Second skin mental pass Elongation Example rate (%)Plating step Note A-a 0.2 Without plating Steel of Invention A-b 0.4Alloyed hot-dip Steel of Invention galvanization B-a 0.3 Without platingSteel of Invention B-b 0.5 Alloyed hot-dip Steel of Inventiongalvanization C-a 0.3 Without plating Steel of Invention C-b 0.5 Hot-dipgalvanization Steel of Invention D-a 1.5 Hot-dip galvanization Steel ofInvention D-b 0.3 Alloyed hot-dip Steel of Invention galvanization E-a0.3 Hot-dip galvanization Steel of Invention E-b 0.5 Alloyed hot-dipSteel of Invention galvanization F-a 0.4 Hot-dip galvanization Steel ofInvention F-b 0.4 Alloyed hot-dip Steel of Invention galvanization G-a0.3 Hot-dip galvanization Steel of Invention G-b 0.3 Alloyed hot-dipSteel of Invention galvanization H-a 0.3 Hot-dip galvanization Steel ofInvention H-b 0.3 Alloyed hot-dip Steel of Invention galvanization I-a0.3 Without plating Steel of Invention I-b 4.5 Alloyed hot-dip Steel ofInvention galvanization J-a 1.8 Without plating Steel of Invention J-b0.3 Alloyed hot-dip Steel of Invention galvanization

TABLE 5 Experi- Second skin mental pass Elongation Example rate (%)Plating step Note K-a 0.3 Without plating Steel of Invention K-b 0.4Alloyed hot-dip Steel of Invention galvanization L-a 2.5 Without platingSteel of Invention L-b 0.3 Alloyed hot-dip Steel of Inventiongalvanization M-a 0.3 Without plating Comparative Steel M-b 0.3 Alloyedhot-dip Comparative Steel galvanization N-a 0.3 Without platingComparative Steel N-b 0.4 Alloyed hot-dip Comparative Steelgalvanization O-a 0.3 Without plating Comparative Steel O-b 0.3 Alloyedhot-dip Comparative Steel galvanization P-a 0.5 Hot-dip galvanizationComparative Steel P-b 0.4 Alloyed hot-dip Comparative Steelgalvanization Q-a 0.3 Hot-dip galvanization Comparative Steel Q-b 0.3Alloyed hot-dip Comparative Steel galvanization R-a 0.3 Without platingComparative Steel R-b 0.3 Alloyed hot-dip Comparative Steelgalvanization S-a 0.4 Without plating Comparative Steel S-b 0.3 Alloyedhot-dip Comparative Steel galvanization T-a 0.3 Without platingComparative Steel T-b 0.4 Alloyed hot-dip Comparative Steelgalvanization

TABLE 6 First skin Annealing Hot rolling pass Maximum Heating FinishingCoiling Elongation heating Holding Left side of Right side ofExperimental Steel temperature temperature Cooling rate temperature ratetemperature time Expression (1) Expression (1) Example No. (° C.) (° C.)(° C./s) (° C.) (%) (° C.) (sec) (° C.) (° C.) A-c A 1100 900 40 450 0.2660 240 68 1026 A-d 1200 890 35 460 0.1 680 200 54 948 A-e 1210 910 40500 0.6 650 250 75 1065 A-f 1230 900 30 510 0.3 790 200 −23 519 A-g 1220910 35 550 0.5 650  20 75 1065 A-h 1230 900 30 580 1.0 680 1210  54 948A-i 1220 890 35 680 0.3 650 300 75 1065 A-j 1210 890 35 630 0.3 680 10054 948 A-k 1220 900 40 550 0.0 720  40 26 792 A-l 1200 910 40 560 0.4660 150 68 1026 A-m 1190 870 45 230 0.7 680 300 54 948 A-n 1210 760 45560 0.6 710 320 33 831 A-o 1210 900 40 470 0.3 660 320 68 1026 B-c B1200 905 45 570 0.5 680 240 54 948 B-d 1210 920 45 650 0.5 700  60 40870 B-e 1220 910 30 500 0.8 520 600 166 1572 B-f 1230 900 35 510 2.5 630600 89 1143

TABLE 7 First skin Annealing Hot rolling pass Maximum Heating FinishingCooling Coiling Elongation heating Holding Left side of Right side ofExperimental Steel temperature temperature rate temperature ratetemperature time Expression Expression Example No. (° C.) (° C.) (°C./s) (° C.) (%) (° C.) (sec) (1) (° C.) (1) (° C.) B-g B 1210 890 35530 2.1 680 1100  54 948 B-h 1220 920 40 550 4.3 610  60 103 1221 B-i1230 930 45 580 6.2 680 200 54 948 B-j 1200 910 30 520 2.2 650 630 751065 B-k 1210 915 45 530 1.0 630 300 89 1143 B-l 1210 920 45 200 0.0 680150 54 948 B-m 1200 910 30 515 0.6 790 300 −23 519 B-n 1210 915 30 5300.5 680  30 54 948 B-o 1220 900 30 550 1.6 640 510 82 1104 C-c C 1200895 45 530 0.5 690 240 47 909 C-d 1210 890 40 430 0.3 700  80 40 870 C-e1230 905 40 490 1.0 680 310 54 948 C-f 1210 910 45 670 1.5 650 500 751065 C-g 1210 915 30 350 0.0 630 800 89 1143 C-h 1220 920 35 515 5.5 660300 68 1026 C-i 1210 890 35 530 2.1 500 300 180 1650

TABLE 8 Experi- Second skin mental pass Elongation Example rate (%)Plating step Note A-c 0.2 Without plating Comparative Steel A-d 0Alloyed hot-dip Steel of Invention galvanization A-e 0.5 Hot-dipgalvanization Steel of Invention A-f 0.1 Alloyed hot-dip ComparativeSteel galvanization A-g 0.5 Alloyed hot-dip Comparative Steelgalvanization A-h 0.3 Alloyed hot-dip Comparative Steel galvanizationA-i 1 Hot-dip galvanization Comparative Steel A-j 1 Hot-dipgalvanization Comparative Steel A-k 0.6 Alloyed hot-dip ComparativeSteel galvanization A-l 2.2 Alloyed hot-dip Steel of Inventiongalvanization A-m 0 Without plating Steel of Invention A-n 0.6 Withoutplating Comparative Steel A-o 0.2 Hot-dip galvanization Steel ofInvention B-c 0.5 Without plating Steel of Invention B-d 0.5 Alloyedhot-dip Comparative Steel galvanization B-e 0.5 Alloyed hot-dipComparative Steel galvanization B-f 0 Without plating Steel of InventionB-g 0.3 Hot-dip galvanization Comparative Steel B-h 0.5 Hot-dipgalvanization Comparative Steel B-i 0.3 Alloyed hot-dip ComparativeSteel galvanization B-j 0.5 Alloyed hot-dip Steel of Inventiongalvanization B-k 0.5 Alloyed hot-dip Steel of Invention galvanizationB-l 0.5 Alloyed hot-dip Comparative Steel galvanization B-m 0.5 Alloyedhot-dip Comparative Steel galvanization B-n 0.3 Without platingComparative Steel B-o 0.3 Alloyed hot-dip Steel of Inventiongalvanization C-c 2.5 Without plating Steel of Invention C-d 0 Hot-dipgalvanization Steel of Invention C-e 1.5 Alloyed hot-dip Steel ofInvention galvanization C-f 0.5 Alloyed hot-dip Comparative Steelgalvanization C-g 0.5 Alloyed hot-dip Comparative Steel galvanizationC-h 0.8 Alloyed hot-dip Comparative Steel galvanization C-i 1 Alloyedhot-dip Comparative Steel galvanization

In Experimental Examples of Tables 2 to 5, the steel sheets wereproduced for the purpose of clarifying the criticalities of the rangesof the component contents of the steel sheets of the present invention.Therefore, the production conditions were set to be in the ranges of thepresent invention. On the other hand, in Experimental Examples of Tables6 to 8, the steel sheets were produced for the purpose of clarifying thecriticalities of the ranges of the production conditions of the presentinvention. Therefore, slabs Nos. A to C were used of which the componentcontents were in the ranges of the present invention.

The properties of the produced steel sheets were evaluated by thefollowing methods.

(Microstructure)

In accordance with the method described in the embodiment, samples weretaken from the portion which was ¼ of the sheet thickness (at a depth of¼ of the sheet thickness) inner from the surface of the steel sheet, andthen the microstructures thereof were observed. Thereafter, themicrostructures were identified, and the area ratio of each structurewas measured by an image analysis method.

The density of Ti(C,N) precipitates and the dislocation density weremeasured by the methods described in the embodiment.

(Tensile Test)

A No. 5 test specimen described in JIS-Z2201 was produced, and a tensiletest was performed in accordance with a test method described inJIS-Z2241. Thereby, the tensile strength (TS), yield strength (yieldstress), and elongation of the steel sheet were measured.

The acceptance range of the elongation depending on the strength levelof the tensile strength was determined by Expression (4) as follows, andthe elongation was evaluated. Specifically, the acceptance range of theelongation was determined in a range of equal to or higher than thevalue of the right side of Expression (4) as follows in consideration ofa balance with the tensile strength.

Elongation [%]≧30−0.02×Tensile Strength [MPa]  (4)

(Hardness)

Using MVK-E micro Vickers hardness tester manufactured by AkashiCorporation, the hardness of a cross-section of the steel sheet wasmeasured. As the hardness (Hvs) of the surface layer of the steel sheet,a hardness at a portion that is 20 μm (at a depth of 20 μm) inner fromthe surface was measured. In addition, as the hardness (Hvc) of thecenter portion of the steel sheet, a hardness at a portion that is ¼ ofthe sheet thickness (at a depth of ¼ of the sheet thickness) inner fromthe surface of the steel sheet was measured. At each portion, hardnessmeasurement was performed three times, and the average of the measuredvalues (average value of n=3) was determined as the hardness (Hvs andHvc). Here, the applied load was set to 50 gf.

(Fatigue Strength and Fatigue Strength Ratio)

The fatigue strength was measured using a Schenck type plane bendingfatigue testing machine in accordance with JIS-Z2275. The stress loadduring measurement was set at a speed of reversed stress testing of 30Hz. In addition, under the above-described conditions, the fatiguestrength was measured at a cycle of 10⁷ by the Schenck type planebending fatigue testing machine. Then, the fatigue strength at the cycleof 10⁷ was divided by the tensile strength measured by theabove-described tensile test; and thereby, a fatigue strength ratio wascalculated. The acceptance range of the fatigue strength ratio was setto be in a range of 0.45 or more.

(Platability)

Platability was evaluated by presence or absence of generation ofnon-plated portions and plating adhesion property.

Whether or not there was a portion which was not plated (a non-platedportion) was visually checked after hot dipping. A steel sheet wherethere was no portion which was not plated was determined as “good(pass)”, and a steel sheet where there is a portion which is not platedwas determined as “bad (fail)”.

In addition, plating adhesion property was evaluated as follows. Aspecimen taken from the plated steel sheet was subjected to a 60 degreesV bending test, and then the specimens on which a bending test wasperformed was subjected to a tape test. In the case where a blackeningof the tape test was less than 20%, the steel sheet was determined as“good (pass)”, and in the case where the blackening of the tape test was20% or more, the steel sheet was determined as “bad (fail)”.

(Chemical Conversion Property)

Using a dip type bond liquid (surface treatment agent) which is commonlyused, the surface of the steel sheet was subjected to a chemicalconversion treatment; and thereby, a phosphate film was formed. Then, acrystalline state of phosphate was observed by a scanning electronmicroscope at 10,000-fold magnification with 5 fields of view. In thecase where crystals of phosphate were precipitated on the entiresurface, the steel sheet was determined as “good (pass)”, and in thecase where there were portions at which crystals of phosphate were notprecipitated was determined as “bad (fail)”.

TABLE 9 Microstructure Mechanical properties Density of Calculated Ti(C, N) Dislocation Yield Tensile result of Experimental Ferrite BainiteMartensite precipitates density stress strength Yield ExpressionElongation Example (%) (%) (%) (/mm³) (/m²) (MPa) (MPa) ratio (4) (%)A-a 85 15 — 2 × 10¹⁰ 2 × 10¹³ 590 640 0.92 17.2 28 A-b 60 40 — — 2 ×10¹³ 570 610 0.93 17.8 26 B-a 30 70 — 1 × 10¹¹ 4 × 10¹³ 760 820 0.9313.6 15 B-b 25 75 — — 4 × 10¹³ 770 830 0.93 13.4 14 C-a 15 85 — — 6 ×10¹³ 915 1010 0.91 9.8 11 C-b 5 70 25 — 6 × 10¹³ 950 1020 0.93 9.6 10D-a 25 75 — — 4 × 10¹³ 790 860 0.92 12.8 13 D-b 20 80 — — 3 × 10¹³ 770850 0.91 13 14 E-a 10 80 10 — 8 × 10¹³ 690 840 0.82 13.2 16 E-b 5 70 25— 8 × 10¹³ 680 830 0.82 13.4 15 F-a 40 60 — — 5 × 10¹³ 590 625 0.94 17.523 F-b 45 55 — — 3 × 10¹³ 570 610 0.93 17.8 22 G-a 30 70 — — 6 × 10¹³770 785 0.98 14.3 18 G-b 35 65 — — 4 × 10¹³ 775 790 0.98 14.2 18 H-a 4060 — 3 × 10¹⁰ 8 × 10¹³ 625 680 0.92 16.4 18 H-b 30 70 — — 6 × 10¹³ 610690 0.88 16.2 19 I-a 10 90 — — 4 × 10¹³ 735 855 0.86 12.9 14 I-b 15 85 —— 4 × 10¹³ 750 840 0.89 13.2 15 J-a 5 70 25 — 4 × 10¹³ 960 995 0.96 10.112 J-b 0 60 40 — 7 × 10¹³ 940 990 0.95 10.2 11

TABLE 10 Microstructure Mechanical properties Density Calculated of Ti(C, N) Dislocation Yield Tensile result of Experimental Ferrite BainiteMartensite precipitates density stress strength Yield ExpressionElongation Example (%) (%) (%) (precipitates/mm³) (/m²) (MPa) (MPa)ratio (4) (%) K-a 30 70 — 2 × 10¹¹ 6 × 10¹³ 810 850 0.95 13 15 K-b 30 6010 — 6 × 10¹³ 830 860 0.97 12.8 14 L-a 0 70 30 — 6 × 10¹³ 960 1120  0.867.6  9 L-b 0 75 25 — 5 × 10¹³ 950 1090  0.87 8.2  9 M-a 90 10 — — 2 ×10¹³ 410 430 0.95 21.4 25 M-b 95 5 — — 1 × 10¹³ 420 440 0.95 21.2 24 N-a0 20 80 — 2 × 10¹⁴ 890 1170  0.76 6.6  7 N-b 0 10 90 — 3 × 10¹⁴ 9001150  0.78 7  7 O-a 50 50 — — 4 × 10¹³ 570 615 0.93 17.7 19 O-b 65 35 —— 3 × 10¹³ 560 620 0.90 17.6 18 P-a 90 10 — — 5 × 10¹³ 440 470 0.94 20.623 P-b 95 5 — — 4 × 10¹³ 430 460 0.93 20.8 22 Q-a 10 80 10 — 7 × 10¹³880 965 0.91 10.7  9 Q-b 5 90  5 — 8 × 10¹³ 890 970 0.92 10.6  8 R-a 4060 — — 7 × 10¹³ 860 930 0.92 11.4 12 R-b 45 55 — — 4 × 10¹³ 870 940 0.9311.2 13 S-a 30 70 — 3 × 10⁸   2 × 10¹³ 580 740 0.78 15.2 19 S-b 20 80 —— 3 × 10¹³ 590 760 0.78 14.8 18 T-a 10 90 — — 9 × 10¹³ 920 990 0.93 10.2 8 T-b 5 95 — — 9 × 10¹³ 910 980 0.93 10.4  8

TABLE 11 Mechanical properties Hardness Hardness of surface of centerHardness Fatigue Fatigue Plating adhesion or Experimental layer portionratio strength strength Chemical conversion Example (Hvs) (Hvc)(Hvs/Hvc) (MPa) ratio properties Note A-a 165 190 0.87 310 0.48 GoodSteel of Invention A-b 160 180 0.89 300 0.49 Good Steel of Invention B-a240 250 0.96 420 0.51 Good Steel of Invention B-b 240 260 0.92 410 0.49Good Steel of Invention C-a 280 300 0.93 460 0.46 Good Steel ofInvention C-b 290 310 0.94 470 0.46 Good Steel of Invention D-a 250 2700.93 400 0.47 Good Steel of Invention D-b 240 260 0.92 390 0.46 GoodSteel of Invention E-a 220 260 0.85 380 0.45 Good Steel of Invention E-b215 250 0.86 380 0.46 Good Steel of Invention F-a 175 190 0.92 320 0.51Good Steel of Invention F-b 170 180 0.94 315 0.52 Good Steel ofInvention G-a 200 230 0.87 370 0.47 Good Steel of Invention G-b 210 2350.89 390 0.49 Good Steel of Invention H-a 200 210 0.95 350 0.51 GoodSteel of Invention H-b 195 215 0.91 340 0.49 Good Steel of Invention I-a215 240 0.90 400 0.47 Good Steel of Invention I-b 220 255 0.86 390 0.46Good Steel of Invention J-a 280 300 0.93 490 0.49 Good Steel ofInvention J-b 270 290 0.93 480 0.48 Good Steel of Invention

TABLE 12 Mechanical properties Plating Hardness Hardness adhesion or ofsurface of center Hardness Fatigue Fatigue Chemical Experimental layerportion ratio strength strength conversion Example (Hvs) (Hvc) (Hvs/Hvc)(MPa) ratio properties Note K-a 260 270 0.96 410 0.48 Good Steel ofInvention K-b 240 260 0.92 420 0.49 Good Steel of Invention L-a 310 3400.91 510 0.46 Good Steel of Invention L-b 290 330 0.88 520 0.48 GoodSteel of Invention M-a 125 130 0.96 205 0.48 Good Insufficient in TSComparative Steel M-b 135 140 0.96 200 0.45 Good Insufficient in TSComparative Steel N-a 260 350 0.74 440 0.38 Good Insufficient in yieldratio, Comparative Steel hardness ratio, and fatigue strength ratio N-b270 340 0.79 460 0.40 Good Insufficient in yield ratio, ComparativeSteel hardness ratio, and fatigue strength ratio O-a 180 190 0.95 3000.49 Bad Deteriorated chemical Comparative Steel conversion propertiesO-b 190 200 0.95 310 0.50 Bad Deteriorated platability Comparative SteelP-a 130 140 0.93 230 0.49 Good Insufficient in TS Comparative Steel P-b140 150 0.93 210 0.46 Good Insufficient in TS Comparative Steel Q-a 270300 0.90 440 0.46 Good Insufficient in elongation Comparative Steel Q-b260 290 0.90 450 0.46 Good Insufficient in elongation Comparative SteelR-a 275 285 0.96 430 0.46 Bad Deteriorated chemical Comparative Steelconversion properties R-b 285 290 0.98 450 0.48 Bad Deterioratedplatability Comparative Steel S-a 175 230 0.76 290 0.39 GoodInsufficient in yield ratio, Comparative Steel hardness ratio, andfatigue strength ratio S-b 170 220 0.77 280 0.37 Good Insufficient inyield ratio, Comparative Steel hardness ratio, and fatigue strengthratio T-a 290 300 0.97 480 0.48 Good Insufficient in elongationComparative Steel T-b 280 290 0.97 470 0.48 Good Insufficient inelongation Comparative Steel

TABLE 13 Microstructure Density of Mechanical properties Ti (C, N)Calculated precipitates Dislocation Yield Tensile result of ExperimentalFerrite Bainite Martensite (precipitates/ density stress strength YieldExpression Elongation Example (%) (%) (%) mm³) (/m²) (MPa) (MPa) ratio(4) (%) A-c 75 25 — — 2 × 10¹³ 400 520 0.77 19.6 26 A-d 75 25 — — 2 ×10¹³ 570 620 0.92 17.6 23 A-e 85 15 —  2 × 10¹¹ 3 × 10¹³ 580 630 0.9217.4 25 A-f 80 20 — 5 × 10⁹ 1 × 10¹³ 520 560 0.93 18.8 25 A-g 90 10 — —1 × 10¹⁴ 510 580 0.88 18.4 25 A-h 90 10 — — 1 × 10¹³ 510 570 0.89 18.624 A-i 98 2 — — 2 × 10¹³ 440 530 0.83 19.4 28 A-j 98 2 — — 2 × 10¹³ 435540 0.81 19.2 27 A-k 90 10 — — 2 × 10¹³ 560 620 0.90 17.6 26 A-l 90 10 —— 3 × 10¹³ 570 610 0.93 17.8 24 A-m 90 10 — — 2 × 10¹³ 580 625 0.93 17.524 A-n 95 5 — — 2 × 10¹³ 500 595 0.84 18.1 25 A-o 80 20 — — 3 × 10¹³ 570630 0.90 17.4 24 B-c 30 70 — — 4 × 10¹³ 730 785 0.93 14.3 18 B-d 35 65 —— 2 × 10¹³ 690 760 0.91 14.8 19 B-e 40 60 — 9 × 10⁹ 3 × 10¹⁴ 700 7600.92 14.8 18 B-f 30 70 — — 4 × 10¹³ 770 820 0.94 13.6 18

TABLE 14 Microstructure Density of Mechanical properties Ti (C, N)Calculated precipitates Dislocation Yield Tensile result of ExperimentalFerrite Bainite Martensite (precipitates/ density stress strength YieldExpression Elongation Example (%) (%) (%) mm³) (/m²) (MPa) (MPa) ratio(4) (%) B-g 20 80 — — 2 × 10¹³ 730 790 0.92 14.2 19 B-h 30 70 — — 2 ×10¹⁴ 720 795 0.91 14.1 18 B-i 30 70 — — 6 × 10¹³ 780 860 0.91 12.8  9B-j 35 65 — — 4 × 10¹³ 720 810 0.89 13.8 18 B-k 30 70 — 2 × 10¹¹ 6 ×10¹³ 730 820 0.89 13.6 18 B-l 30 50 20 — 4 × 10¹³ 680 810 0.84 13.8 19B-m 35 65 — — 4 × 10¹³ 600 760 0.79 14.8 20 B-n 25 75 — — 2 × 10¹⁴ 670780 0.86 14.4 18 B-o 30 70 — — 4 × 10¹³ 730 810 0.90 13.8 18 C-c 20 80 —— 8 × 10¹³ 915 1020 0.90 9.6 12 C-d 10 90 — — 7 × 10¹³ 930 1010 0.92 9.811 C-e 15 85 — — 7 × 10¹³ 920 1015 0.91 9.7 11 C-f 50 50 — — 5 × 10¹³760 960 0.79 10.8 14 C-g 5 50 45 — 9 × 10¹³ 910 1020 0.89 9.6 12 C-h 1090 — — 9 × 10¹³ 970 1105 0.88 7.9  6 C-i 15 85 — — 3 × 10¹⁴ 800 965 0.8310.7 13

TABLE 15 Mechanical properties Plating Hardness Hardness adhesion or ofsurface of center Hardness Fatigue Fatigue Chemical Experimental layerportion ratio strength strength conversion Example (Hvs) (Hvc) (Hvs/Hvc)(MPa) ratio properties Note A-c 130 160 0.81 230 0.44 Good Insufficientin TS, yield ratio, Comparative Steel hardness ratio, and fatiguestrength ratio A-d 160 180 0.89 290 0.47 Good Steel of Invention A-e 170190 0.89 300 0.48 Good Steel of Invention A-f 140 170 0.82 240 0.43 GoodInsufficient in TS, hardness Comparative Steel ratio, and fatiguestrength ratio A-g 150 180 0.83 230 0.40 Good Insufficient in TS,hardness Comparative Steel ratio, and fatigue strength ratio A-h 145 1800.81 235 0.41 Good Insufficient in TS, hardness Comparative Steel ratio,and fatigue strength ratio A-i 135 165 0.82 220 0.42 Good Insufficientin TS, hardness Comparative Steel ratio, and fatigue strength ratio A-j140 170 0.82 230 0.43 Good Insufficient in TS, hardness ComparativeSteel ratio, and fatigue strength ratio A-k 150 190 0.79 260 0.42 GoodInsufficient in hardness ratio Comparative Steel and fatigue strengthratio A-l 175 190 0.92 280 0.46 Good Steel of Invention A-m 180 190 0.95290 0.46 Good Steel of Invention A-n 140 180 0.78 240 0.40 GoodInsufficient in hardness ratio Comparative Steel and fatigue strengthratio A-o 165 185 0.89 295 0.47 Good Steel of Invention B-c 200 230 0.87370 0.47 Good Steel of Invention B-d 180 230 0.78 330 0.43 GoodInsufficient in hardness ratio Comparative Steel and fatigue strengthratio B-e 180 220 0.82 330 0.43 Good Insufficient in hardness ratioComparative Steel and fatigue strength ratio B-f 230 245 0.94 380 0.46Good Steel of Invention

TABLE 16 Mechanical properties Plating Hardness Hardness adhesion or ofsurface of center Hardness Fatigue Fatigue Chemical Experimental layerportion ratio strength strength conversion Example (Hvs) (Hvc) (Hvs/Hvc)(MPa) ratio properties Note B-g 200 240 0.83 340 0.43 Good Insufficientin hardness ratio Comparative Steel and fatigue strength ratio B-h 190240 0.79 330 0.42 Good Insufficient in hardness ratio Comparative Steeland fatigue strength ratio B-i 200 245 0.82 330 0.38 Good Insufficientin elongation, Comparative Steel hardness ratio, and fatigue strengthratio B-j 230 260 0.88 400 0.49 Good Steel of Invention B-k 230 255 0.90390 0.48 Good Steel of Invention B-l 190 250 0.76 330 0.41 GoodInsufficient in hardness ratio Comparative Steel and fatigue strengthratio B-m 170 230 0.74 310 0.41 Good Insufficient in yield ratio,Comparative Steel hardness ratio, and fatigue strength ratio B-n 175 2400.73 320 0.41 Good Insufficient in hardness ratio Comparative Steel andfatigue strength ratio B-o 225 260 0.87 390 0.48 Good Steel of InventionC-c 270 310 0.87 470 0.46 Good Steel of Invention C-d 265 305 0.87 4650.46 Good Steel of Invention C-e 265 305 0.87 470 0.46 Good Steel ofInvention C-f 250 300 0.83 380 0.40 Good Insufficient in yield ratio,Comparative Steel hardness ratio, and fatigue strength ratio C-g 240 3100.77 390 0.38 Good Insufficient in hardness Comparative Steel ratio andfatigue strength ratio C-h 280 340 0.82 370 0.33 Good Insufficient inelongation, Comparative Steel hardness ratio, and fatigue strength ratioC-i 230 300 0.77 360 0.37 Good Insufficient in hardness ComparativeSteel ratio and fatigue strength ratio

At first, the influences of the components of the steel materials aredescribed.

The C amounts of steels Nos. M and N are out of the range of the presentinvention. The steel sheets (Experimental Examples M-a and M-b) producedusing the steel No. M were insufficient in strength. The steel sheets(Experimental Examples N-a and N-b) produced using the steel No. N wereinsufficient in yield ratio and fatigue strength ratio.

The Si amounts and Al amounts of steels Nos. O and R were greater thanthe ranges of the present invention. The steel sheets (ExperimentalExamples O-a, O-b, R-a, and R-b) produced using the steels Nos. O and Rhad problems with plating adhesion property and chemical conversionproperty.

The Mn amounts of steels Nos. P and Q are out of the range of thepresent invention. The steel sheets (Experimental Examples P-a and P-b)produced using the steel No. P were insufficient in strength. The steelsheets (Experimental Examples Q-a and Q-b) produced using the steel No.Q were insufficient in elongation.

The Ti amounts of steels Nos. S and T are out of the range of thepresent invention. The steel sheets (Experimental Examples S-a and S-b)produced using the steel No. S were insufficient in yield ratio andfatigue strength ratio. The steel sheets (Experimental Examples T-a andT-b) produced using the steel No. T were insufficient in elongation.

Next, the influences of the production conditions are described.

In Experimental Example A-c, the heating temperature of the slab duringhot rolling was insufficient; and thereby, TiC could not be dissolved inaustenite. Therefore, the produced steel sheet was insufficient instrength and fatigue strength.

In Experimental Example A-n, the finishing temperature during hotrolling was reduced. Therefore, the produced steel sheet wasinsufficient in fatigue strength ratio.

In Experimental Examples A-i, A-j, B-d, and C-f, since the coilingtemperatures during hot rolling were high, amounts of solid-solubilizedTi (solid-solution Ti) in the hot rolling stage became insufficient.Therefore, the produced steel sheets were insufficient in fatiguestrength ratio.

In Experimental Examples A-k, B-l, and C-g, since the elongation ratesof the first skin pass rolling after the hot rolling were insufficient,introduction of strains to the surface layers of the steel sheets becameinsufficient. As a result, the precipitation effect in the surface layerafter annealing was not sufficiently obtained. Therefore, the producedsteel sheets were insufficient in fatigue strength ratio.

In Experimental Examples B-i and C-h, since the elongation rates of thefirst skin pass rolling after the hot rolling were excessively high, theinfluence of the processing strains was increased. Therefore, theproduced steel sheets were insufficient in elongation and fatiguestrength ratio.

In Experimental Examples A-f and B-m, since the annealing temperaturesafter the first skin pass rolling were high, precipitates coarsened.Therefore, fatigue strength ratios and densities of precipitates of theproduced steel sheets were degraded.

In Experimental Examples B-e and C-i, since the annealing temperaturesafter the first skin pass rolling were low, precipitation of TiC did notsufficiently proceed. Therefore, the produced steel sheets wereinsufficient in fatigue strength ratio.

In Experimental Examples A-g, B-h, and B-m, since the holding times in atemperature range of 600° C. or higher during the annealing after thefirst skin pass rolling were short, precipitation of TiC did not proceedsufficiently. Therefore, the produced steel sheets were insufficient infatigue strength ratio.

In Experimental Examples A-h and B-g, since the holding times in atemperature range of 600° C. or higher during the annealing after thefirst skin pass rolling were long, precipitates coarsened. Therefore,the produced steel sheets were insufficient in fatigue strength ratio.

The microstructures of the steel sheet of the present invention(Experimental Example B-k) and the comparative steel (ExperimentalExample B-e) were compared to each other. In the steel sheet of thepresent invention (Experimental Example B-k), precipitation of TiCoccurred during annealing, and as shown in FIGS. 11 and 13, the densityof precipitates having sizes of 10 nm or smaller was increased to1.82×10¹¹ precipitates/mm³. In contrast, in the comparative steel sheet(Experimental Example B-e), precipitation of TiC did not proceed asdescribed above, and as shown in FIGS. 12 and 14, the density ofprecipitates having sizes of 10 nm or smaller was maintained at about8.73×10⁹ precipitates/mm³.

INDUSTRIAL APPLICABILITY

In accordance with the present invention, a high-strength steel sheet, ahot-dipped steel sheet, and an alloyed hot-dipped steel sheet can beprovided which have a tensile strength in a range of 590 MPa or more andwhich are excellent in fatigue properties, elongation and collisionproperties. In the case where they are applied to components for anautomobile, a reduction in the weight and enhancement of safety of theautomobile can be achieved. In particular, the hot-dipped steel sheetand the alloyed hot-dipped steel sheet of the present invention have theabove-described excellent properties and excellent rust prevention.Therefore, they can be applied to chassis frames, and they cancontribute to the reduction in the weight of an automobile. As describedabove, the present invention can be appropriately applied to fields ofsteel sheets for automobile components such as chassis frames.

1-7. (canceled)
 8. A method for producing a high-strength steel sheethaving excellent fatigue properties, elongation and collision propertiescomprising: heating a slab comprising: in terms of percent by mass %,0.03 to 0.10% of C; 0.01 to 1.5% of Si; 1.0 to 2.5% of Mn; 0.1% or lessof P; 0.02% or less of S; 0.01 to 1.2% of Al; 0.06 to 0.15% of Ti; and0.01% or less of N; and containing as the balance, iron and inevitableimpurities, at a temperature in a range of 1,150 to 1,280° C. andperforming hot rolling under conditions where a finish rolling isfinished at a temperature in a range of not less than an Ar₃ point,thereby obtaining a hot-rolled material; coiling the hot-rolled materialin a temperature range of 600° C. or less, thereby obtaining ahot-rolled steel sheet; subjecting the hot-rolled steel sheet to acidpickling; subjecting the pickled hot-rolled steel sheet to first skinpass rolling at an elongation rate in a range of 0.1 to 5.0%; annealingthe hot-rolled steel sheet under conditions where a maximum heatingtemperature (Tmax° C.) is in a range of 600 to 750° C. and a holdingtime (t seconds) in a temperature range of 600° C. or higher fulfillsexpressions (1) and (2) as follows; and subjecting the annealedhot-rolled steel sheet to second skin pass rolling,530−0.7×Tmax≦t≦3,600−3.9×Tmax  (1)t>0  (2).
 9. The method for producing the high-strength steel sheethaving excellent fatigue properties, elongation and collision propertiesaccording to claim 8, wherein an elongation rate is set to be in a rangeof 0.2 to 2.0% in the second skin pass rolling.
 10. The method forproducing the high-strength steel sheet having excellent fatigueproperties, elongation and collision properties according to claim 8,wherein ½ or more of the amount of Ti contained in the hot-rolled steelsheet after the coiling exists in a solid-solution state.
 11. A methodfor producing a hot-dipped steel sheet having excellent fatigueproperties, elongation and collision properties comprising: heating aslab comprising: in terms of percent by mass %, 0.03 to 0.10% of C; 0.01to 1.5% of Si; 1.0 to 2.5% of Mn; 0.1% or less of P; 0.02% or less of S;0.01 to 1.2% of Al; 0.06 to 0.15% of Ti; and 0.01% or less of N; andcontaining as the balance, iron and inevitable impurities, at atemperature in a range of 1,150 to 1,280° C. and performing hot rollingunder conditions where a finish rolling is finished at a temperature ina range of not less than an Ar₃ point, thereby obtaining a hot-rolledmaterial; coiling the hot-rolled material in a temperature range of 600°C. or less, thereby obtaining a hot-rolled steel sheet; subjecting thehot-rolled steel sheet to acid pickling; subjecting the pickledhot-rolled steel sheet to first skin pass rolling at an elongation ratein a range of 0.1 to 5.0%; annealing the hot-rolled steel sheet underconditions where a maximum heating temperature (Tmax° C.) is in a rangeof 600 to 750° C. and a holding time (t seconds) in a temperature rangeof 600° C. or higher fulfills expressions (1) and (2) as follows, andperforming hot dipping to form a hot-dipped layer on a surface of thehot-rolled steel sheet, thereby obtaining a hot-dipped steel sheet; andsubjecting the hot-dipped steel sheet to second skin pass rolling,530−0.7×Tmax≦t≦3,600−3.9×Tmax  (1)t>0  (2).
 12. The method for producing the hot-dipped steel sheet havingexcellent fatigue properties, elongation and collision propertiesaccording to claim 11, wherein an elongation rate is set to be in arange of 0.2 to 2.0% in the second skin pass rolling.
 13. A method forproducing an alloyed hot-dipped steel sheet having excellent fatigueproperties, elongation and collision properties comprising: heating aslab comprising: in terms of percent by mass %, 0.03 to 0.10% of C; 0.01to 1.5% of Si; 1.0 to 2.5% of Mn; 0.1% or less of P; 0.02% or less of S;0.01 to 1.2% of Al; 0.06 to 0.15% of Ti; and 0.01% or less of N; andcontaining as the balance, iron and inevitable impurities, at atemperature in a range of 1,150 to 1,280° C. and performing hot rollingunder conditions where a finish rolling is finished at a temperature ina range of not less than an Ar₃ point, thereby obtaining a hot-rolledmaterial; coiling the hot-rolled material in a temperature range of 600°C. or less, thereby obtaining a hot-rolled steel sheet; subjecting thehot-rolled steel sheet to acid pickling; subjecting the pickledhot-rolled steel sheet to first skin pass rolling at an elongation ratein a range of 0.1 to 5.0%; annealing the hot-rolled steel sheet underconditions where a maximum heating temperature (Tmax° C.) is in a rangeof 600 to 750° C. and a holding time (t seconds) in a temperature rangeof 600° C. or higher fulfills expressions (1) and (2) as follows,performing hot dipping to form a hot-dipped layer on a surface of thehot-rolled steel sheet so as to obtain a hot-dipped steel sheet, andsubjecting the hot-dipped steel sheet to an alloying treatment toconvert the hot-dipped layer into an alloyed hot-dipped layer; andsubjecting the hot-dipped steel sheet on which the alloying treatment isperformed to second skin pass rolling,530−0.7×Tmax≦t≦3,600−3.9×Tmax  (1)t>0  (2).
 14. The method for producing the alloyed hot-dipped steelsheet having excellent fatigue properties, elongation and collisionproperties according to claim 13, wherein an elongation rate is set tobe in a range of 0.2 to 2.0% in the second skin pass rolling.